text:
\section*{\sffamily \normalsize  INTRODUCTION}
The understanding of the microscopic mechanisms underlying the plastic response of amorphous solids to externally driven deformations
is a current issue in material science research
both for the lack of a complete theoretical background and its importance in technical applications. \cite{ArgonFractures,StachurskiProgPolymSci97,BarratShearReview,KramerQuestions}
Solids subjected to small deformations respond linearly as expected from elasticity theories. \cite{ArgonPlasticTheory,EshelbyElastic,BudianskyElastic,WardElastic}
An increasing strain on the system causes the increase of internal stress. { Focusing on pure shear deformation,} the elastic modulus $G$ of the system under the studied deformation can be derived
from the slope of the stress-strain curve in the small strain regime \cite{BarratShearReview}
both locally and globally \cite{BarratAQS13}.
Upon increasing strain, amorphous solids show complex and far from linear behavior \cite{BarratPlastic02,BarratPlastic04,BarratPlastic05},
with heterogeneous and  protocol-dependent \cite{ProcacciaProtocol,ProcacciaCooling} phenomena taking place mainly due to the absence of long-range order \cite{ArgonFractures}.
{ Having reached a characteristic yield strain, corresponding to the shear yield strength $\tau_Y$, the transition from the (reversible) elastic state to the (irreversible) plastic
one takes place \cite{StachurskiProgPolymSci97,RobbinsHoyJPolymSci06,RottlerSoftMatter10}.
In an ideal elasto-plastic body (Hooke-St.Venant) $\tau_Y$ is the maximum stress \cite{StachurskiProgPolymSci97}.}

{
Despite the complexity of the plastic behavior in amorphous solids at the local scale, some general features have been found in the macroscopic quantities. An interesting aspect of yielding
is that the yield stress is proportional to the elastic modulus. In particular,
for a linear, dislocation-free array of atoms Frenkel derived long time ago the relation   $\tau_Y / G\simeq 1/ ( \pi \sqrt{3} )  \simeq 0.18 \, $ at $T= 0 K$
\cite{FrenkelYield26,FrenkelYieldModel,StachurskiProgPolymSci97}. A more recent experimental work found $\tau_Y/ G  \sim 0.11$ for polymers \cite{ArgonFractures} and the universal value
$0.036 \pm 0.002$  for metallic atomic glasses \cite{JohnsonMetallic}. The ratio $\tau_Y/ G$ depends on the temperature and, for a given temperature, is universal for metallic glasses up to slightly
below the glass transition temperature \cite{JohnsonMetallic}. The finding has been interpreted in terms of similar inter-particle potentials \cite{LernerProcaccia}.
The microscopic origin of the proportionality between  $\tau_Y$ and $G$ has been rationalized by noting that in both metals and polymers the yield stress is primarily governed by energy storing elastic
processes: dislocation line energy in metals, strain energy around molecular kinks in polymers \cite{ArgonPlasticTheory}. Since the elastic modulus in glassy polymers is dominated by
inter-molecular forces  \cite{BerniniElasticJPolB15}, it was concluded that the energy barriers to plastic flow in glassy polymers were dominated by intermolecular rather than intra-molecular
interactions \cite{ArgonFractures}}, so some similarities
can be found in the comparison between atomic and polymeric systems \cite{ArgonSilicon}.
On the other hand, the intra-molecular interactions can have a primary role in determining the structure of a polymer solid upon cooling from the liquid phase, which is of great importance
to determine the elastic properties of the final structure \cite{ProcacciaProtocol, ProcacciaCooling}.

\begin{figure}[t]
\begin{center}
\includegraphics[width= \linewidth]{fig1.eps}
\end{center}
\caption{{ \small Illustrative snapshots of the different polymer solids under study: {semicrystalline polymer} (A), disordered {glassy polymer} (B),
nematic {glassy polymer} (C). Monomer position is emphasized in the
left column, bond orientation is emphasized in the right column. Differently from the {nematic glassy polymer},
both {semicrystalline and disordered glassy polymers} exhibit local bond disorder. The snapshots refer to
chains with bond length $l_b=1.12$ and bending stiffness $k_\theta= 0$ (A), $4$ (B), $12.5$ (C).}}
\label{MDbox}
\end{figure}

The aim of this work is to investigate the existence of the correlation of shear elastic modulus $G$ and yielding stress $\tau_Y$ in polymer solids by means of molecular dynamics (MD) simulations.
A model is presented in which the systematic variation of characteristic parameters of the intra-molecular interactions, namely the bond length $l_b$  and the bending stiffness $k_\theta$ of pairs of
contiguous bonds  in a chain, leads to different {semicrystalline, disordered or nematic} structures, see \cref{MDbox}.
{ For fully-flexible chains with no bending potential ( $k_\theta= 0$ ) we find that the
yield stress increases with the elastic modulus in a way which is very close to the universal law of the atomic metallic glasses \cite{JohnsonMetallic}, suggesting that, in the absence of
bending stiffness, connectivity and structure play minor roles in the yield process of the present polymer model.}
Increasing the bending stiffness of the chains causes the increasing growth of the local nematic ordering of near chains.
It is seen that the onset of nematic order {\it increases} the elastic modulus $G$ and {\it  decreases} the yielding stress $\tau_Y$, thus evidencing the different influence of the local order
on the plasticity and the elasticity. A tentative explanation of the subsequent reduction of the ratio $\tau_Y/G$ in terms of  the distributions of the per-monomer stress is offered.

\section*{ \sffamily \normalsize NUMERICAL METHODS}
\label{numerical}
We consider a coarse-grained polymer model of $N_c=160$ linear, unentangled chains with $M=25$ monomers per chain.
The total number of monomers is $N=4000$. Non-bonded monomers at distance $r$ interact via the truncated and shifted Lennard-Jones (LJ) potential:
\begin{equation}
U^{LJ}(r)=\varepsilon\left [ \left (\frac{\sigma^*}{r}\right)^{12 } - 2\left (\frac{\sigma^*}{r}\right)^6 \right]+U_{cut}
\end{equation}
for $r\leq r_c=2.5\sigma$ and zero otherwise, where $\sigma^{\ast}=2^{1/6}\sigma$, is the position of the potential minimum with depth $\varepsilon$.
The value of the constant $U_{cut}$ is chosen to ensure that $U_{LJ}\left(r\right)$ is continuous at $r=r_c$. Henceforth, all quantities are expressed in terms of reduced units:
lengths in units of $\sigma$, temperatures in units of $\varepsilon/k_B$ (with $k_B$ the Boltzmann constant) and time $\tau_{MD}$ in units of $\sigma \sqrt{m / \varepsilon}$ where
$m$ is the monomer mass. We set $m = k_B = 1$. The bonding interaction is approximated via the harmonic potential
\begin{equation}
U_{bond}(r)=k_b(r-l_b)^2
\end{equation}
where $l_b$ is the equilibrium bond length and $k_b=300\varepsilon/\sigma^2$ is the bond rigidity. Differently from previous studies  concerning fully-flexible chains \cite{UnivSoftMatter11,PuosiLepoJCPCor12,PuosiLepoJCPCor12_Erratum,Puosi11}, the bending angle interaction between adjacent chemical bonds is included through a potential of the form \cite{KarayiannisBendingAngle15}:
\begin{equation}
U_{bending}=k_\theta(1-\cos\theta_b)
\end{equation}
where $k_\theta$ is the bending stiffness, $\cos\theta_b^i=\vec{b}_{i+1}\cdot\vec{b}_i/||\vec{b}_{i+1}|||\vec{b}_i||$ and the bond vector $\vec{b}_i=\vec{r}_{i+1}-\vec{r}_i$, where $\vec{r}_i$
is the position of the i-th monomer.
Periodic boundary conditions are used. The study was performed in the $NPT$ ensemble (constant number of particles, pressure and temperature).
The integration time step is set to $\Delta t=0.005$ time units \cite{Puosi11,UnivPhilMag11}.
The simulations were carried out using LAMMPS molecular dynamics software (http://lammps.sandia.gov). \cite{PlimptonLAMMPS}

A systematic study is performed by changing the bond length $l_b$ and bending stiffness $k_\theta$. We focus on two families of systems: fully flexible polymers
($k_\theta=0$) with $0.91\leq l_b\leq1.12$, and semi-flexible/stiff polymers with $l_b=1.12$ and $1.0\leq k_\theta\leq12.5$.
All samples are equilibrated in the $NPT$ ensemble at $P=0$. They are initially equilibrated at the following temperatures: $T=1.2$ for $k_\theta < 7$, $T=1.4$ for
$ 7 \le k_\theta < 12.5$, $T=1.6$ for $k_\theta = 12.5$. Then, they are isobarically cooled down to $T=0$ with a constant quench rate of $|\dot{T}|=2\cdot10^{-6}$.
Both the equilibration and the quench procedures are close to the one adopted in Ref. \cite{KarayiannisBendingAngle15}.
 Isobaric quenches have also been considered  in other MD investigations of plastic yield in glassy polymers. \cite{RottlerSoftMatter10}
After the quench, simple shear deformations of the polymer solids at $T=0$, $P=0$ are performed via the Athermal Quasi-Static (AQS) protocol described in details in Ref. \cite{BarratShearReview}.
Initially, the undeformed simulation box containing the sample is a cube with side $L$.
An infinitesimal strain increment $\Delta\varepsilon=10^{-5}  L$ is applied, after which the system is allowed to relax
in the local potential energy minimum via a suitable minimization algorithm. The procedure is repeated up to the total strain $\varepsilon_{tot}=15 \cdot 10^{-2} L$.

Simple shear is performed independently in the planes ($xy$, $xz$, $yz$), and at each strain step in the plane {$\alpha\beta$ the corresponding component of the macroscopic stress tensor
$\tau_{\alpha, \beta}$ is taken as the average value of the per-monomer stress  $\tau_{\alpha, \beta}^i$:
\begin{equation}
\label{stresstensor1}
\tau_{\alpha, \beta} =   \frac{1}{N}  \sum_{i=1}^N  \tau_{\alpha, \beta}^i
\end{equation}
In an athermal system the expression of the per-monomer stress in the atomic representation is \cite{allenMolPhys}:
\begin{equation}
\label{stresstensor2}
\tau_{\alpha, \beta}^i =  \frac{1}{2 \, v} \sum_{j \ne i} r_{\alpha ij} F_{\beta ij}
\end{equation}
where $F_{\gamma  k l}$ and $r_{\gamma  k l}$  are the $\gamma$ components of the force between the $k$th and the $l$th monomer and their  separation, respectively, and $v$ is the average per-monomer
volume, i.e. $v = L^3/N$.}
For each plane we then obtain a stress-strain curve, an illustrative example of which is given in \cref{carico}. The result is quite analogous to what reported for many other systems under
athermal conditions \cite{MottArgonSuter93,FalkViscoPlastic,Maeda2Dplastic,MelandroPlastic,LemaitrePlastic,ProcacciaPiecewise}  with an initial linear increase followed by increasing bending and onset
of the plastic regime. In particular, similarly to other MD studies of glassy polymers \cite{YieldMDPolymerMM15}, one notices that, in the plastic regime,  the stress levels off to a plateau with
fluctuations caused by subsequent loading phases and sudden stress drops. We point out that the initial non-zero stress in the unstrained solid seen in  \cref{carico} is a well-known phenomenon
usually ascribed to the limited size of the simulation cell \cite{HawardCristPhysGlassyPolym}.

We measure the shear elastic modulus $G$ as the slope of the stress-strain curve in the linear regime ($\varepsilon \le 0.01$), via the relation $G  = \tau / 2\varepsilon $, see inset of \cref{carico}.
Following Ref. \cite{LernerProcaccia}, the yield stress $\tau_Y$ is taken as the average value of the stress after the first significant plastic event, defined as the first stress drop
of at least $\Delta\tau_{th}=0.1$, see \cref{carico}. { This choice is  consistent with other definitions in the presence \cite{RottlerSoftMatter10}, or not \cite{RobbinsHoyJPolymSci06},
of strain softening, i.e. the reduction in stress following yield}. The results are robust with respect to changes of $\Delta\tau_{th}$.
Data concerning 16 distinct simulation runs are gathered for each physical state. Each run is averaged over the three planes $xy$, $xz$ and $yz$.

 %
\begin{figure}[t]
\begin{center}
\includegraphics[width=0.95\linewidth]{fig2.eps}
\end{center}
\caption{ { \small Typical stress-strain curve of our polymer solids under athermal, quasi-static, simple-shear deformation. After a first 'loading' phase, a   plateau-like plastic regime sets
in where a series of sudden stress drops are observed. The yield stress $\tau_Y$ is defined as the average
value of $\tau$ in the plastic regime  \cite{LernerProcaccia}. The elastic modulus $G$ (see inset) is measured via a linear fit of the stress-strain curve in the linear regime of small deformations
$2\epsilon<0.02$. The plot refers to a system of fully-flexible chains ($k_\theta=0$) with bond length $l_b =1.12$.}}
\label{carico}
\end{figure}

\section*{\sffamily \normalsize RESULTS AND DISCUSSION}
\label{resultsdiscussion}

\subsection*{ \sffamily \normalsize Structural analysis during quench-cooling}
\label{quench}

The elastic properties of amorphous solids strongly depend on the sample preparation \cite{ProcacciaCooling,ProcacciaProtocol}.
Thus, we preliminarily characterize the most relevant structural changes of our systems occurring during the isobaric quench from the liquid  to the athermal solid.

In order to study more rigorously the structural order of the systems, we resort to the order parameters defined by Steinhardt \textit{et al.} \cite{Steinhardt83}.
One considers in a given coordinate system the polar and azimuthal angles $\theta({\bf r}_{ij})$ and $\phi({\bf r}_{ij})$ of the
vector ${\bf r}_{ij}$ joining the $i$-th central monomer with the $j$-th one belonging to the neighbors within a
preset cutoff distance $r_{cut} = 1.2 \; \sigma^* \simeq 1.35$ \cite{Steinhardt83}. $r_{cut}$ is a convenient definition of
the first coordination shell size \cite{sim}. The vector ${\bf r}_{ij}$ is usually referred to as a ``bond'' and has
not to be confused with the {\it actual} chemical bonds of the polymeric chain.
To define a global measure of the order in the system, one then introduces the quantity:
\begin{equation} \label{Qbarlm_glob}
	\bar{Q}_{lm}^{glob}=\frac{1}{N_{b}}\sum_{i=1}^{N}
\sum_{j=1}^{n_b(i)}Y_{lm}\left[\theta({\bf
r}_{ij}),\phi({\bf r}_{ij})\right]
\end{equation}
where $n_b(i)$ is the number of bonds of $i$-th particle, $N$ is the total number of particles in the system, $Y_{lm}$
denotes a spherical harmonic and $N_b$ is the total number of bonds:
\begin{equation} \label{N_b}
	N_b=\sum_{i=1}^{N} n_b(i)
\end{equation}
The global orientational order parameter $Q_l^{glob}$ is defined by \cite{Steinhardt83}:
\begin{equation} \label{Ql_glob}
 Q_l^{glob}=\left [ \frac{4\pi}{(2l+1)} \sum_{m=-l}^{l}
|\bar{Q}_{lm}^{glob}|^2 \right ]^{1/2}
\end{equation}
The above quantity is invariant under rotations of the coordinate system and takes characteristic values which can be used to quantify the kind and the degree of
rotational symmetry in the system \cite{Steinhardt83}. In the absence of {\it large}-scale order, the bond orientation is uniformly distributed around the unit sphere and $Q_l^{glob}$
is rather small since it vanishes as $ \sim N_b^{-1/2}$ \cite{RintoulTorquatoJCP96}. On the other hand, $Q_6^{glob}$ is very sensitive to any kind of crystallization and increases significantly
when order appears \cite{GervoisGeometrical99,GiuntoliCristallo}.
A local orientational parameter $Q^{loc}_{l}$ can also be defined. We define the auxiliary quantity
\begin{equation} \label{Qbarlm_loc}
	\bar{Q}^{loc}_{lm}( i )=\frac{1}{n_b(i)}\sum_
{j=1}^{n_b(i)}Y_{lm}\left[\theta({\bf r}_{ij}),\phi({\bf
r}_{ij})\right]
\end{equation}
The local order parameter $Q^{loc}_{l}$ is defined as \cite{Steinhardt83}:
\begin{equation} \label{Ql_loc}
 Q^{loc}_{l}=\frac{1}{N} \sum_{i=1}^{N}  \left [
\frac{4\pi}{(2l+1)} \sum_{m=-l}^{l} |\bar{Q}^{loc}_{lm}( i
)|^2 \right ]^{1/2}
\end{equation}
In general $Q^{loc}_{l}\ge Q_l^{glob}$. In the presence of ideal order, {\it all} the particles have the {\it same} neighborhood configuration, and the equality $Q^{loc}_{l} = Q_l^{glob}$
follows.

We first examine the density and the global order of fully-flexible chains ($k_\theta=0$).
The global positional order of the monomers is monitored via the Steinhardt order parameter $Q^{glob}_6$. \cref{BNDdensq6} plots the increase of both the density $\rho$ and
the order parameter $Q^{glob}_6$ for different bond lengths $l_b$ while decreasing the temperature at constant pressure $P=0$ from the initial liquid state to the final solid state.
Fully flexible polymers either exhibit global order or glassify upon cooling, depending on the bond length $l_b$.
Global order is revealed by sharp jumps in density $\rho$ and $Q^{glob}_6$ for $l_b=1.06,1.09,1.12$.
A local-order analysis, presented later in the paper, clarifies that the states with global order are {semicrystalline polymers} with coexisting polymorphs at $T=0$.
Systems with shorter bond length form {glassy polymers}, with no significant global order.  If the bond length is comparable to the monomer size, $l_b \approx \sigma^\ast \simeq 1.12$,
the formation of ordered structures is to be expected \cite{SimmonsBcc13,KarayiannisCEE09,KarayiannisPRL09HSC,KarayiannisSM10AthHSC,GiuntoliCristallo}, whereas shorter bonds are known
\cite{Hamley2007,DoyeFrustratCrystalPhysChemChemPhys07,MilsteinPRB70,SimmonsBcc13}
to cause geometrical frustration which hinders the crystallization process, thus favoring the formation of {disordered glassy polymers}.

\begin{figure}[t]
\begin{center}
\includegraphics[width=  \linewidth]{fig3}
\end{center}
\caption{ { \small Density $\rho$ (upper panel) and global order parameter $Q^{glob}_6$ (lower panel) of a melt of fully-flexible chains ($k_\theta=0$) with different bond length
$l_b$ during the isobaric quench
from the liquid to the solid phase. Chains with short bond length ($l_b\leq1.03$) form {disordered glassy polymers} since the bond length is  incommensurate with the
Lennard-Jones length scale $\sigma^\ast \simeq 1.12$.
Chains with bond length comparable to $\sigma^\ast \simeq 1.12$ exhibit steep increase of the density $\rho$ and the global order parameter $Q^{glob}_6$ upon cooling.
 In the latter case, the local-order analysis presented in Fig.\ref{corrq6q4} clarifies that the corresponding solids at $T=0$ are {semicrystalline polymers}
with coexisting polymorphs.}}
\label{BNDdensq6}
\end{figure}

We now turn to semi-flexible and stiff chains ($k_\theta > 0$). Since the reduced flexibility favors local nematic ordering,
i.e. the alignment of near bonds, we divide the sample in $n^3$ cells with side $L/n$ and  define the bond-orientation order parameter in the $i$-th cell as \cite{LuoSommerLocalNematicOrderM11}
\begin{equation}
{S_i}=\sqrt{\frac{3}{2}Tr(q_i^2)}, \quad q_{i,\alpha\beta}=\langle\hat{b}_{\alpha}\hat{b}_{\beta}-\frac{1}{3}\delta_{\alpha\beta}\rangle_i
\label{orderparameter0}
\end{equation}
where $1\le i\le n^3$, $Tr$ is the trace operator,  $q_i$ is a $3 \times 3$ orientational tensor with components of $q_{i,\alpha\beta}$,
$\hat{b}_{\alpha}$ and $\hat{b}_{\beta}$ are the Cartesian components of the normalized bond vectors $\vec{b}$ and the statistical average $\langle...\rangle_i$ is performed
on all the bonds of the $i$-th cell. {Following Karayiannis and coworkers \cite{KarayiannisBendingAngle15},
we initially choose $n=6$ corresponding to cells with side of about $2-3$ monomer diameters.
An average local bond-orientation order parameter is then defined as \cite{LuoSommerLocalNematicOrderM11}
\begin{equation}
S_{loc}^{bond}= \frac{1}{216} \sum_{i=1}^{216} S_{i}
\label{orderparameter}
\end{equation}
}
\begin{figure}[t]
\begin{center}
\includegraphics[ width= 0.85 \linewidth]{fig4.eps}
\end{center}
\caption{{ \small Density $\rho$ (top), global order parameter $Q^{glob}_6$ (middle) and local bond-orientation order parameter {$S_{loc}^{bond}$}
(bottom) of a melt of chains with increasing
bending stiffness during the isobaric quench from the liquid to the solid phase. Bond length $l_b=1.12$. For  fully-flexible chains ($k_\theta=0$) a steep increase of the density $\rho$ and the global
order parameter $Q^{glob}_6$ is revealed at $T\simeq 0.5$. For semi rigid/stiff chains ($k_\theta>0$): i) all the final solid states are {glassy polymers} ($Q^{glob}_6 < 0.05$), ii)
on cooling, the bending stiffness triggers a transition to a nematic state with bond ordering occurring in the liquid phase and freezing below the glass transition.}}
\label{STFdensS}
\end{figure}

\begin{figure}[t]
\begin{center}
\includegraphics[ width=   \linewidth]{fig5.eps}
\end{center}
\caption{{ \small Dependence of the {global} bond-orientation order parameter {$S_{glob}^{bond}$} and the global {chain-orientation} order parameter
{$S_{glob}^{Ree}$} on the bending stiffness by increasing $k_\theta$ at $T=0$.
Bond length $l_b=1.12$. Note that, even for high bending stiffness, the global order is not strong despite the local ordering shown in Fig.\ref{STFdensS} (bottom). See Fig.\ref{MDbox} for selected snapshots.}}
\label{S1Sg}
\end{figure}

{The $S_{loc}^{bond}$ order parameter  ranges between $S_{loc}^{bond}=1$ (perfect alignment) and $S_{loc}^{bond}=0$ (random orientation).}
\cref{STFdensS} plots the density $\rho$, the order parameter $Q^{glob}_6$ and
the local bond-orientation order parameter {$S_{loc}^{bond}$} of systems with bond length $l_b=1.12$ and different bending stiffness  $k_{\theta}$, during the isobaric quench
from the initial liquid state to the final solid state.
The latter exhibits global order only if the chains are fully flexible ($k_{\theta}=0$), as signaled by the jumps of both the density and the global order
parameter at $T\simeq 0.5$, otherwise {glassy polymers} with small global order ($Q^{glob}_6 < 0.05$) are obtained. It is seen that the increasing bending stiffness of the chains
triggers a transition to a nematic state with considerable local alignment of the bonds, as detected by the increase of the bond-orientational order parameter {$S_{loc}^{bond}$}.
The resulting local orientational order freezes below the glass transition, yielding a {nematic glassy polymer}.

It is interesting to consider the {global} bond-orientation order. To this aim, we set $n=1$ and define the {global} bond-orientation order parameter $S_{glob}^{bond}$ as $S_1$ from eq.\ref{orderparameter0} to perform the average of the bond orientation over a {\it single} cell coinciding with all the sample.
The quantity is plotted in Fig.\ref{S1Sg}. On increasing the  bending stiffness $k_\theta$ at $T=0$, {$S_{glob}^{bond}$} starts from $\sim 0.05$ for fully-flexible chains ($k_\theta=0$),
then increases and levels off at the plateau level $S_{glob}^{bond} \simeq 0.38$ for  $k_\theta \gtrsim 8.5$. This suggests that the sample is {\it locally oriented}
(high {$S_{loc}^{bond}$}), but {\it macroscopically nearly isotropic} (small {$S_{glob}^{bond}$}) for strong bending stiffness.
To corroborate the previous conclusion, we consider the alignment of the end-to-end unit vector of the chains via the global {chain-orientation} order parameter
{$S_{glob}^{Ree}$} \cite{KarayiannisBendingAngle15}.
By construction, {$S_{glob}^{Ree}$} spans the range between {$S_{glob}^{Ree} = 1$} (perfect alignment of all the chains) and
{$S_{glob}^{Ree} = 0$} (random orientation). Fig.\ref{S1Sg} shows that {$S_{glob}^{Ree}$} increases with the bending stiffness but it is not large.

\begin{figure}[t]
\begin{center}
\includegraphics[width= \linewidth]{fig6.eps}
\end{center}
\caption{{ \small Bivariate distributions of the local order parameters $Q^{loc}_{4}$ and $Q^{loc}_{6}$ for characteristic states at $T=0$:  (a) {semicrystalline polymer} with
$l_b=1.12$ and $k_\theta=0$;
(b) {semicrystalline polymer} with $l_b=1.06$ and $k_\theta=0$; (c) disordered {glassy polymer} with $l_b=1.12$ and $k_\theta=4$;
(d) nematic {glassy polymer} with $l_b=1.12$ and $k_\theta=12.5$. In panel a) the regions corresponding
to the BCC, FCC and HCP structures at the level of the first neighbor shell are marked. The region "D" labels states with first neighbor shell different from the BCC, FCC and HCP ones.
The contour lines have equal contour interval and divide the whole elevation range evenly.
}}
\label{corrq6q4}
\end{figure}

In order to gain more insight into the structure of the polymeric solids Fig.\ref{corrq6q4} presents the correlation plots of the local order parameters $Q^{loc}_{4}$ and $Q^{loc}_{6}$,
characterizing the order of the {\it first} neighbor shell of each monomer. Fig.\ref{corrq6q4}a shows the complex nature of the solid state with $l_b=1.12$ and $k_\theta = 0$, corresponding to
fully-flexible chains with bond length comparable to the monomer size.
Four different regions with highly-correlated pairs
$(Q^{loc}_{4},Q^{loc}_{6})$  are apparent. According to previous studies \cite{GiuntoliCristallo,LocalOrderJCP13}, two of them signal face-centered cubic (FCC) and hexagonal close packed (HCP)
local packings. For the same polymer model with $k_\theta = 0$ and $l_b=1.12$ at  $T=0$, FCC and HCP close packed structures together with other (unspecified) non close-packed environments were detected
\cite{KarayiannisBendingAngle15}.
We also identify high correlations in the region $(Q^{loc}_{4},Q^{loc}_{6}) \simeq (0.12,0.4)$. These values are ascribed to a deformed body-centered cubic (BCC) structure
with $(Q^{loc}_{4},Q^{loc}_{6})$ pair significantly different from the ideal BCC { due to poor stability of the BCC lattice \cite{GiuntoliCristallo,MisraCrystStability1940}}.
On the basis of previous studies \cite{FrenkelLJCrystalJCP96}, we believe that such BCC structures were nucleated as metastable regions during the quench and frozen in the solid phase at $T=0$.
BCC structures have been reported for the present model with $k_\theta = 0$ and $l_b \simeq 0.97$ in the crystallization of a polymer melt exposed to well-ordered walls \cite{SimmonsBcc13} and in
the spontaneous isothermal crystallization of an unbounded polymer melt \cite{GiuntoliCristallo}.  { The D region in Fig.\ref{corrq6q4}a represents environments with first neighbor shell different
from the BCC, FCC and HCP ones. }

In summary, the solid state of fully-flexible chains with bond length comparable to the monomer size, $l_b =1.12  \approx \sigma^\ast$,  is semicrystalline with coexisting polymorphs.
The structure of the solid appears to be much less heterogeneous by decreasing the bond length or increasing the bending stiffness. In fact, Fig.\ref{corrq6q4}b shows that,
if $l_b = 1.06$ with $k_\theta = 0$,  the D region is enhanced to the detriment of the BCC, FCC and HCP regions. For  $l_b \le 1.03$ and $k_\theta = 0$ the solid is a disordered glass and only the
D region is apparent (not shown). A similar finding is observed by keeping $l_b =1.12$ and increasing the strength of the bending potential, see c) and d) panels of Fig.\ref{corrq6q4}.
{ Then, we see that the D region is characteristic of our glassy systems}.

We note that Fig.\ref{corrq6q4}d shows two weak lobes located at $Q^{loc}_{6} \simeq 0.48$ with $Q^{loc}_{4} \simeq 0.09$ and $0.175$. By comparison with panels a) and c) of Fig.\ref{corrq6q4},
the finding suggests
reentrant FCC and HCP ordering on increasing the strength of the bending potential with $l_b =1.12$. The finding is consistent with the results reported by Karayiannis and coworkers
\cite{KarayiannisBendingAngle15} where the fraction of sites with close-packed order (FCC or HCP similarities) is close to one in systems with
{$S_{glob}^{Ree} \simeq 1$}, i.e. nearly straight chains,
and high local orientation order, {$S_{loc}^{bond} \sim 0.95$}. We remind that in our case {$S_{glob}^{Ree}$} and {$S_{loc}^{bond}$} are not larger than about $0.33$ and $0.7$,
respectively. Incidentally, the fact that we find less
global and local orientational order with the same polymer model with respect to Ref.\cite{KarayiannisBendingAngle15}  is ascribed to the smaller size of our sample.

\begin{figure}[t]
\begin{center}
\includegraphics[width=0.99 \linewidth]{fig7.eps}
\end{center}
\caption{ { \small Correlation plot of  the elastic modulus $G$ and the average yield stress $\tau_Y$ for the athermal solids made by fully-flexible chains with different bond lengths.
The error bars of $\tau_Y$ and $G$ reflect the fluctuation of the stress during the steady state of the plastic regime, and the uncertainty of the fit in the linear elastic regime, see \cref{carico},
respectively.  {Semicrystalline polymers} and disordered glassy polymers exhibit correlations in two different regions, inside of which the influence of the bond length is minor.
{The black continuous line is the
universal law of metallic glasses $\tau_Y = m \, G$ with slope $m = 0.036 \pm 0.002$ \cite{JohnsonMetallic}. The uncertainty on the $m$ parameter is bounded by the two dashed lines.}}
}
\label{BNDcorr}
\end{figure}

\subsection*{\sffamily \normalsize Correlation between yield stress and shear modulus }
\label{shear}
\cref{BNDcorr} is a correlation plot of the average yield stress $\tau_Y$ and the elastic shear modulus $G$  for the solids made by fully-flexible chains with different bond lengths $l_b$.
The plot presents the data on a run-by-run basis, i.e. no average between runs with the same bond length is performed.
A general tendency of the yield stress $\tau_Y$ to increase with the modulus is observed.
It is seen that {disordered glassy polymers} exhibit limited changes of both $\tau_Y$ and $G$, whereas
 {semicrystalline polymers} show a wider distribution across the different runs. We ascribe the effect to the polymorphic character of the ordered solids \cite{GiuntoliCristallo}.
Also, {semicrystalline polymers} show higher $G$ and $\tau_Y$ values with respect to {disordered glassy polymers}, meaning that the increased order of the monomeric
arrangement causes the system to react to shear deformations with stronger internal stresses with respect to its amorphous counterpart both in the linear regime and at the yield point.
In \cref{BNDcorr} we superimpose to our data the characteristic universal law of the metallic {\it atomic} glasses, i.e. the line $\tau_Y = m \, G$ with  $m = 0.036 \pm 0.002$ \cite{JohnsonMetallic}.
Deviations are apparent but not large, thus suggesting that, in the absence of bending stiffness, connectivity and structure play minor roles in the yield process of the present polymer model.

\begin{figure}[t]
\begin{center}
\includegraphics[width=0.99 \linewidth]{fig8.eps}
\end{center}
\caption{{\small
Correlation plot of  the elastic modulus $G$ and the average yield stress $\tau_Y$ for  the athermal solids made by chains with different bending stiffness.
Bond length $l_b=1.12$. As in \cref{BNDcorr} the black continuous line is the universal law of metallic glasses $\tau_Y = m \, G$ with slope $m = 0.036 \pm 0.002$ \cite{JohnsonMetallic}.
Differently from {semicrystalline and disordered glassy polymers}, {\it {nematic glassy polymers}} exhibit large deviations from that law.}}
\label{STFcorr}
\end{figure}

The introduction of bending stiffness, $k_{\theta}\neq0$, and the subsequent nematic order provide a different scenario. This is clearly visible in the correlation plot of the average yield stress
$\tau_Y$ and the elastic shear modulus $G$, see \cref{STFcorr}. For low and intermediate bending stiffness, $k_{\theta}\le 4$, the solids are {semicrystalline polymers}
or microscopically disordered {glassy polymers}
respectively, with ratio  $\tau_Y / \, G$  close to the characteristic universal value $0.036 \pm 0.002$ of the atomic metallic glasses \cite{JohnsonMetallic}, as in \cref{BNDcorr}.
For {nematic glassy polymers}, $k_{\theta}\ge 7$, the ratio $\tau_Y / \, G$ decreases by increasing the bending stiffness of the chain.

\begin{figure}[t]
\begin{center}
\includegraphics[width=0.99\linewidth]{fig9.eps}
\end{center}
\caption{Correlation between the elastic modulus (top) and the yield stress $\tau_Y$ (bottom) with the  local bond-orientation order parameter {$S_{loc}^{bond}$,
Eq.\ref{orderparameter}.}
Bond length $l_b=1.12$. The dashed line divides the regions pertaining to disordered (left) and nematic (right) {glassy polymers}.
It is seen that in {nematic glassy polymers} the increasing local bond alignment {\it increases} the elastic modulus and {\it decreases} the yield stress.}
\label{decoupling}
\end{figure}
\begin{figure}[t]
\begin{center}
\includegraphics[width=0.99\linewidth]{fig10.eps}
\end{center}
\caption{Correlation between the elastic modulus (top) and the yield stress $\tau_Y$ (bottom) with the density. Bond length $l_b=1.12$.
The dashed line divides the regions pertaining to disordered (left) and nematic (right) {glassy polymers}.
Notice that the {nematic glassy polymers} with $k_\theta = 7$ and $k_\theta = 8.5$ have
nearly identical densities but rather different local nematic order, see \cref{STFdensS} top. The fact that their modulus and yield stress are appreciably different
signals the influence of the bond ordering.}
\label{decoupling2}
\end{figure}

We have investigated the origin of the deviations of the ratio $\tau_Y / \, G$ from the characteristic universal value of the atomic metallic glasses.  Elasticity and yielding of polymeric solids
are both affected by density \cite{StachurskiProgPolymSci97} and local nematic order
\cite{DouglasThermoset02,DouglasNematicThermosetMM98,OrtizNematic}, two properties which are changed  by varying the  bending stiffness, see \cref{STFdensS}.
We first consider the influence of nematic order.  \cref{decoupling} shows that in {nematic glassy polymers}, on increasing the local orientational  order of the bonds,
the elastic modulus {\it increases} and the yield stress {\it decreases}.
A similar effect has been observed by Ortiz et al \cite{OrtizNematic} in the glassy phase of a macroscopically disordered, liquid-crystalline thermoset,
where changing the densely cross-linked network structure from an ensemble of randomly oriented rigid-rods to local nematic increases the modulus and decreases the yield stress,
see Table 3 and 4 of Ref. \cite{OrtizNematic}.  Since the increase of the nematic order is accompanied by the decrease of the density (apart from a small inversion on increasing $k_{\theta}$ from
7 to 8.5, see \cref{STFdensS} top), we have also examined the role of the density. \cref{decoupling2}
shows that in disordered {glassy polymers}, in spite of a density change of about $6 \%$ neither $G$ nor $\tau_Y$ change appreciably.
Changes are visible in {nematic glassy polymers} where density changes are smaller due to the better packing. This suggests that density plays a minor role, with respect to nematic order,
in setting both the modulus and the yield stress. In this regard, the comparison between the {nematic glassy polymers} with bending stiffness $k_{\theta}=7$ and $k_{\theta}= 8.5$
provides more insight.
The two systems have rather comparable density but quite different local nematic order, see \cref{STFdensS}. \cref{decoupling2} shows that their moduli (yield stress) are distinctly different,
increasing (decreasing) with the local nematic order. All in all, the discussion of \cref{decoupling} and \cref{decoupling2} points to the conclusion that in the  polymer model under study elasticity
and yielding are more affected by the local nematic order than packing. The weak role of packing was also noted in other studies concerning the fast dynamics of polymers  \cite{BarbieriGoriniPRE04}.

{Finally, \cref{distrib} plots the { per-monomer} shear stress distributions  in {semicrystalline polymers} ($k_\theta = 0$), disordered ($k_\theta = 4.0$) and
nematic ($k_\theta = 12.5$) {glassy polymers}.
It is seen that the nematic {glassy polymer} exhibits the broadest distribution with heavy non-gaussian tails. This finding suggests a tentative explanation
of the reduction of the ratio $\tau_Y/G$ in {nematic glassy polymers} with respect to {semicrystalline polymers} and
disordered {glassy polymers}, see \cref{STFcorr}.
In fact, it is known that application of a local stress $\tau'$ decreases the energy barrier $\Delta E$ for plastic rearrangements to $\Delta E - \tau' V^\star$ where $V^\star$ is
an activation volume \cite{StachurskiProgPolymSci97,ArgonPlasticTheory}. If the energy barrier is due to the elastic resistance of the surroundings  treated as an isotropic continuum,
one finds  $\Delta E = G V^\dagger$, where $V^\dagger$ is a further activation volume distinct from $V^\star$  \cite{Frenkel,StachurskiProgPolymSci97,LiGilmanElsticDisclinJAppPhys70,ArgonFractures}.
If one assumes that yielding at $T= 0K$ occurs when the energy barrier vanishes, one finds that the local yield  stress $\tau'_Y = V^\dagger/ V^\star G$. If the stress distribution is narrow,
the local stresses little differ from the average stress,  $\tau_Y \simeq \tau'_Y = V^\dagger/ V^\star G$, and one recovers the usual coupling between the elastic modulus and the macroscopic yield
stress. Otherwise, if the distribution broadens, highly stressed regions yield when the average stress is much less than their stress $\tau'$, so that $\tau_Y < V^\dagger/ V^\star G$, namely the ratio
$\tau_Y/ G$ decreases with respect to the characteristic value for systems with narrow stress distribution.  We are aware that our arguments are rather rough. Nonetheless, they offer a consistent
picture leading to the scenario of \cref{STFcorr}.}

\begin{figure}[t]
\begin{center}
\includegraphics[width=0.99\linewidth]{fig11.eps}
\end{center}
\caption{Distribution of the { per-monomer} shear stress, Eq.\ref{stresstensor2}, in {semicrystalline polymers} ($k_\theta = 0$), disordered ($k_\theta = 4.0$) and nematic
($k_\theta = 12.5$) {glassy polymers}.
All systems have  bond length $l_b=1.12$ and subjected to average stress $\tau = 0.5$, exceeding the linear elastic regime but still far from the region where the sharp plastic drops are observed,
see \cref{carico}. Inset: comparison between the distributions of the disordered  and the {nematic glassy polymers}.
The dotted curves are the best-fit with gaussians showing that the distribution of the nematic {glassy polymer} exhibits heavy non-gaussian tails.}
\label{distrib}
\end{figure}

\section*{\sffamily \normalsize CONCLUSIONS}
\label{conclusions}
Elasticity and yielding in polymer solids have been investigated by MD simulations of a coarse-grained model of linear chains with different bending stiffness and bond length. Following the isobaric quench at $T=0$, three kind of distinct structures are observed:
\begin{itemize}
\item {disordered glassy polymers}: systems with no positional order and and no bond-orientational order,
\item {nematic glassy polymers}: systems with no local positional order but a strong degree of local bond ordering,
\item semicrystalline polymers: systems with local positional order and no bond-orientational order.
\end{itemize}
Note that in this model system semicrystalline polymers do not have any bond orientational order but in other models, e.g. the CG-PVA model \cite{MeyerMullerPlatheMM2001}, short chains with $M \le 30$ form unfolded semicrystalline structures with both local positional (2D hexagonal) and local bond-orientational order \cite{MeyerMullerPlatheJCP2001,MeyerFold02}.

Under simple shear deformations, it is found that in systems with bond disorder the ratio $\tau_Y/G$ between the
shear yield strength $\tau_Y$ and the shear modulus $G$ is close to the universal value of the atomic metallic glasses. In the presence of increasing nematic ordering the shear modulus of the
{glassy polymer}
increases while the shear yield strength decreases, thus reducing the ratio $\tau_Y/G$. The finding parallels similar experimental results concerning nematic thermosets. The results suggest that
nematic order has stronger influence than density on elasticity and yielding. { A tentative explanation of the reduction of the ratio $\tau_Y/G$ in {nematic glassy polymers} with
respect to {semicrystalline polymers} and disordered {glassy polymers} is offered, pointing out the larger width of the { per-monomer} stress distributions.}

\subsection*{\sffamily \normalsize ACKNOWLEDGMENTS}

F. Puosi is gratefully thanked for helpful discussions. A generous grant of computing time from IT Center, University of Pisa and Dell EMC${}^\circledR$ Italia is gratefully acknowledged.

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\end{document}

redpajama_set_name: RedPajamaArXiv


I've got a tag plugin that needs to operate on the deformed points of an object.
What is the best way of retrieving this information through the C++ API? It seems like GetDeformCache() doesn't always work, depending on where and when you call it. The documents briefly mention using the Hierarchy class to rebuild caches, but state that this is extremely resource intensive and should be used sparingly.
Since I'm using the tag on anywhere from 10 to 80 objects in a single scene, I'm guessing I probably shouldn't be using the Hierarchy class as per the SDK warning. However, I'm not sure what else I should use in it's place. I need to get access to the deformed points whenever the tag is executed, but as I said before, GetDeformCache() tends to fail if the scene has just been loaded from disk or is being sent to an external renderer (like the picture viewer).
There's gotta be some way of doing this, since the Xpresso Point node has a "Use deformed points" option that does exactly what I'm looking for. I just need to know how to implement that in my own C++ plugin.
deformers have been evaluated, therefore you will have access to the deformed geometry.
This seems to work reasonably well most of the time, but sometimes it fails when I'm rendering to the picture viewer or render queue. I haven't been able to track down the exact cause, GetDeformCache() just returns nothing, even though I know there should be active deformations (which means the stuff the tag positions onto the geometry isn't in the right place during render).
Do you by chance know how the Xpresso "Point" node handles deformation internally? Are they just forcing the cache build via SendModelingCommand() with CSTO or the Hierarchy class?
can you tell us a bit more about where and how you use GetDeformCache()? Perhaps some simplified code example?
The only code I have right now is the former version of the plugin I wrote under Python, since I'm attempting to port it to C++.
In the Python plugin, I was trying to get the deform cache for the object under the Execute() method. I would only attempt to get the deform cache if the target object had children objects underneath it (which I realize is poor design, since I'm assuming that the child objects are deformers). If I failed to get the deform cache on the first try, I'd call SendModelingCommand() with CSTO on that object, which has the undocumented side effect of forcing a cache rebuild for that object (yes, I know the documentation states I shouldn't do this for that exact reason). Then I'd try and get the deform cache again, and give up if that failed.
This was, of course, all a really horrible kludgy hack.
The actual plugin itself reads in geometry from a deformed polygon object, and creates a set of splines or hair guides (depending on what object the tag is attached to) travelling across the surface of a quadrangle polygon strip. I'm basically using this to approximate the functionality of "Geo Maya Hair", which is a Mays plugin that turns polygon strips into hair guides.
I want to get this thing ported to C++. My primary problem is that I don't know what the correct method is to get the deformed points of a polygon object **under any circumstances**. Ideally, I think the plugin should check GetDeformCache() first just to see if the cache is available, but if it isn't, then I need to do something else to manually rebuild the cache. I don't know how to do that, and the hacks I used in the Python version (as detailed above) certainly don't seem like the right way to go about that.
I tried here with a small example, basically a TagData plugin running on Generator priority, where I use the code snippet on GetCache(). It seems to work as expected. For example I get caches when rendering to Picture Viewer or on scene load. But then again, my test scene might be lacking the needed complexity. Can you maybe provide me with a scenario, where you have problems retrieving a cache? Or maybe you can send me a scene file?
My plugin was originally designed to operate in conjunction with a hair object. It generates custom hair guides based on strips of quadrangle polygon geometry. It also has provisions for selecting a specific "strip" of polygons from a polygon object containing numerous separated strips.
Maybe I was just configuring the priorities wrong, but I could never get things working reliably through GetDeformCache(). Sooner or later, it'd always return nil/none/null, which would screw up the renders somewhere (usually in the PV or render queue). Which is when I came up with that workaround for forcing the cache building.
How can I force a C++ plugin build the deform caches for an object if they have not already been built?
I'm not aware of a possibility to start building the caches from a TagData plugin.
I think, it would be better to solve your problems retrieving the caches correctly. But as said before I will need your help, as I have not been able to reproduce any issues.
I am not 100% sure if this will actually work but you could try calling the CurrentStateToObject modelling command. There is actually a wrapper method that does this in lib_sculpt.h.
Personally I'd rather like to find the reason, why caches can not be retrieved the way it is intended. Such workarounds might not be the best solution performance-wise.
I've gotten enough of the C++ plugin written that I could start fiddling around with it in a variety of test scenes.
It seems to work fine. The priority controls on the tag definitely have an effect, since if I set them to anything lower then "Generators", I can see GetDeformCache() fail and return nothing. If I set the priority to "Generators" or higher, then GetDeformCache() returns a valid object on the first try.
Sorry for the confusion, everything works as intended.
The Python version of the plugin definitely *does* have issues with this, but the C++ version does not. I just kinda assumed the C++ version would share the same issues as the Python plugin, and I'd need some way of ensuring that GetDeformCache() returned a valid object. Looks like that's not the case.


* Push to device was aborted.
Service provisioning failed while trying to provision storage.
I successfully re-discovered Vblock to confirm all usernames, passwords were OK and ECOM service is running but still could not provision storage.
I also recently upgraded one VNX array from 5.32.000.5.207 to 5.32.000.5.209 but it was not the issue as the other array on 5.32.000.5.207 exhibited the same errors.


text:
Cette page concerne les évènements survenus en 1942 en Grèce  :

Évènement
 État grec, régime fantoche et collaborateur mis en place par l'occupant nazi et fasciste (1941-1944)
 Grande Famine (1941-1944)
 juin : Opération Albumen
  : Rafle de la place Eleftherías à Thessalonique.
  :
  : Konstantínos Logothetópoulos devient le deuxième Premier ministre du gouvernement collaborateur, succédant à Georgios Tsolakoglou.

Création
 décembre : Armée populaire de libération nationale grecque
 Libération nationale et sociale

Dissolution
 Organisation patriote socialiste hellénique
 Comité suprême de la lutte crétoise

Sport
  (football)

Naissance
 Dimitri Bertsekas, mathématicien appliqué, ingénieur électricien et informaticien.
, diplomate.
 Isaac Kapetas, rabbin.
Dimítrios Kremastinós, député.
Tónia Marketáki, journaliste, réalisatrice, scénariste, actrice, productrice et monteuse de films.
Lilí Mavrokefálou, poétesse.
 , actrice.

Décès
 Nikólaos Balános, architecte.
, écrivain.
 , peintre.
 Geórgios Polymenákos, général, commandant-en-chef de l'armée lors de la campagne d'Asie mineure.

Articles connexes
 Histoire de la Grèce pendant la Seconde Guerre mondiale
 Occupation de la Grèce pendant la Seconde Guerre mondiale
 Histoire de la Grèce

Chronologie de la Grèce au XXe siècle

redpajama_set_name: RedPajamaWikipedia


G.R.A.M Podcast Volume 32 presented by LSB (Soul R, Spearhead).
DJ System-D takes to the controls for a summarizing mix of poster Drum & Bass forthe 26th Festival de Dour in Balzaal.
Loxy is a Don and this set recorded live in Brooklyn April 26th, 2014 at Reconstrvct XVIII is the cleanest, darkest set that purely demonstrates just what this mater of DNB is all about.
The Kongkast podcast is back this week with one of their favorite duos in the bass music scene – Andy Hobbs and Dan Blishen aka Silent Dust. The duo has been making some serious waves in the various BPM movement such as their unique cross genre between Jungle and Juke.
Beagle-808 brings us some new deep drum & bass for the summer, well for the next couple of days at least. Enjoy!


TODO DA is a multi-disciplinary design and architecture firm based in New York City led by Jorge Prado, RA, LEED AP and Silva Ajemian. Founded in 2005, TODO DA specializes in residential, commercial, and furniture design. The firm has extensive experience with apartment, loft and townhouse renovations, new houses, extensions to existing buildings, offices, showrooms, restaurants, and schools in and around NYC.
Prior to founding TODO DA, Jorge worked on adaptive reuse projects from Asia Society and FIAF to fashion house Elie Tahari's office facilities in Millburn, NJ. Silva worked on a number of NYC Branch Libraries and conversions of industrial buildings to residential condominiums.
The firm's approach to designing a project is to collaborate with clients to achieve spaces that reflect their needs and desires. These collaborations enrich the creative process. The end result is a uniquely tailored material and spatial environment. Jorge and Silva enjoy working with constraints and turning them into opportunities. Each project is a provocation that yields a fine attention to detail and execution through the use of the latest technologies and environmentally friendly materials where possible.
The firm has been recognized with awards and publications internationally. They have participated in exhibitions, design conferences and trade fairs in Palermo, Tokyo and New York. Jorge and Silva have taught the New Jersey Institute of Technology. They are two of the founding members of First Street Green, a community based arts organization working with the New York City Department of Parks and Recreation, Partnership for Parks and arts organizations citywide to revitalize and to bring art programming to First Park on the lower east side of Manhattan.


A quick introduction: I work for NAG as a Technical Support Engineer in Germany. I look after clients in Germany, Switzerland and Austria and am interested in approximation and interpolation methods as well as .NET languages. Naturally, I'm big fan of the NAG library for .NET as well as the NAG Library for SMP and Multicore.
As a non FORTRAN programmer to be able to run some tests with the new NAG Library for SMP and Multicore I decided to choose a few simple tests that I could run using both CPUs of my laptop. I created some .NET console applications, wrappers around the FORTRAN routines that were able to pass all the arguments to the routines in a .NET fashion and implemented a time counting function. There are many parts of the Library parallelized, including quadrature, partial differential equations, interpolation, curve and surface fitting, linear algebra, correlation and regression analysis, multivariate methods, random number generators, time series analysis, sorting and special functions. I chose routines from the correlation, curve and surface fitting and random number generation chapters.
The results were calculated on a Intel® Core™2 Duo Prozessor P8700 (2.53GHz,1066MHz,3MB) machine with Windows Vista as operating system and Microsoft Visual Studio 2008 (32‐bit project) as compiler. Each function was used to solve a large enough problem to allow for parallelism. Each test case was run with one or two threads. The number of threads was set using the system variable 'OMP_NUM_THREADS' in the Visual Studio console window.
After compiling the examples with the C# compiler option 'csc', the number of threads is first set to the value 1 and then set to the value 2. The only result which is printed is the time which the routine needs for calculating the result.
NAG routine C06PKF calculates the circular convolution or correlation of two complex vectors of period n.
In this example to complex vectors are build up with a period of n (n=5000000) and the correlation and the circular convolution are calculated.
NAG routine G05YKF generates a quasi-random sequence from a log-normal distribution. It must be preceded by a call to one of the initialization routines G05YLF or G05YNF. The number N (N = 10000000) of quasi-random numbers of dimension IDIM (IDIM = 4) are passed and also the mean (IDIM) of the underlying Normal distribution for each dimension and the std (IDIM) standard deviation of the underlying Normal distribution.
NAG Routine E02CAF forms an approximation to the weighted, least-squares Chebyshev series surface fit to data arbitrarily distributed on lines parallel to one independent coordinate axis. It determines a bivariate polynomial approximation of degree k in x and l in y to the set of data points , with weights , for and . That is, the data points are on lines , but the x values may be different on each line. The polynomial is represented in double Chebyshev series form.
L, the required degree of y was set also to 100.
Also the fitting polynomial was evaluated at the data points using the routine E02CBF.
Others were excited to learn and understand they might use their second processor!
Never show a C# program to HPC Computer Programmers who love Fortran and think languages such as C are radical!
NAG Library for SMP and Multicore really is easy to use and for those who are current NAG Fortran Library users the transition really is painless. Users just need to take care of the number of threads used.
I should admit a bias as a NAG employee, but perhaps more importantly I'd like to use this blog to talk to users / prospective users and offer help. As I said in my introduction I am interesting in this, but also hearing from those who have interests in approximation and interpolation. I plan to work with some of my developer colleagues to expand some of NAG's interpolation solvers.
My article shows how the NAG SMP library can be used with .NET in exactly the same manner as the NAG serial library. The performance of the SMP library is achieved by careful tuning of the NAG source code using OpenMP to parallelise the code. For the PC environment the SMP library offered uses the Intel Fortran compiler and hence issues Intel OpenMP threads.
Because of this, the SMP library works best in this threading environment or when called from a single thread. Use from a different threading environment can cause difficulties. It is worth noting that my examples were timed and issued from a single .NET thread.
If you are planning to use the NAG SMP library from a multi-threaded .NET program then it we recommend that the environment variable OMP_NUM_THREADS be set to one. This switches off the threading in the SMP library and avoids thread conflict with the .NET threads of the main program.


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Where to Find Bigfoot in the Pacific Northwest
If you're going to catch a Sasquatch, you have to know where to look.
By Allison Williams 6/28/2016 at 11:45am
Big feet, big snowboard at Espresso Chalet on Highway 2.
Image: Joe Schumacher
Where, exactly, is Bigfoot? If he's real—and our peek into the secret lives of Bigfoot hunters proves that plenty of people think he is—then the wild and woody Pacific Northwest is a likely home for a giant man-ape. We pinpointed six destinations for serious Bigfoot hunting. If see him, send us a photo.
Oregon's (humane) Bigfoot trap
1. Bigfoot's Trap
Venture west of Medford, Oregon, to the Collings Mountain Trail at Applegate Lake, then hike three-quarters of a mile near an abandoned miner's cabin. That small wooden structure with a bolted-open gate? It's a Bigfoot trap, the world's first and only. Built in 1974 on an old mining claim by industrious researchers, it's now maintained by the Forest Service and covered in graffiti. Back when it was in service, it only ever caught a few bears. Maybe it needs a bigger block of cheese bait.
2. Bigfoot's Country Retreat
Why wouldn't Bigfoot hang out in the Blue Mountains, timber-covered hills and grasslands near Walla Walla? After all, it's close to wine country. Forest Service worker Paul Freeman claimed to have filmed a Sasquatch in the area in 1994, and his reports grabbed the attention of Washington State University cryptozoologist Grover Krantz. Bigfoot sightings are unverified, but we do know that the world's largest organism lives in the Oregon wing of the Blue Mountains: the massive Armillaria ostoyae fungus.
Bigfoot on the Spirit Lake Highway: such a poser.
3. Bigfoot's Grave
When Mount St. Helens blew in 1980, it definitely killed 57 people and, say believers, it might've taken out Bigfoot. About halfway up the Spirit Lake Highway, the route to the mountain's Johnston Ridge Observatory, a 28-foot Sasquatch statue stands guard at the North Fork Survivors Gift Shop. The smiling figure is a cheerful contrast to the somber memorials to known St. Helens destruction, like a buried A-frame house and the site of vulcanologist Johnston's final research.
4. Bigfoot's Coffee Fix
Seminal 1980s psychodrama Harry and the Hendersons (just kidding, it's a John Lithgow family comedy the New York Times called "aggressively awful") was filmed around Index, and the Highway 2 Espresso Chalet has plenty of memorabilia to mark its brush with Hollywood. One Sasquatch statue holds a snowboard, clearly hoping to bum a ride to Stevens Pass from anyone stopping for Bigfoot cookies and lattes.
Footloose at Sasquatch Days
Image: Tourism Harrison
5. Bigfoot's Hot Tub
The Sasquatch name comes from the British Columbia Sts'ailes tribe's word for "wild man," so the small Canadian town of Harrison Hot Springs celebrates the myth every year with Sasquatch Days (this year June 25-26). The fest is a unique combo of First Nation cultural events and Bigfoot hunts, but expeditions are available year-round through Sasquatch Country Adventures. Located just to the north: Sasquatch Provincial Park. Even when the big guy doesn't show his hairy face, the lakeside village more regularly offers hot springs soaks and a killer prime rib dinner.
6. Bigfoot's Home Base
Bluff Creek, California, is the Bigfoot world's Roswell, its Bermuda Triangle. It's the site of the most famous Sasquatch footage ever captured—or the site of the biggest Bigfoot hoax ever perpetuated. The Patterson-Gimlin film is still debated by cryptozoologists almost 50 years after it was made, but the surrounding Humboldt County woods have been known for Bigfoot encounters for even longer. Should he prove elusive, the Willow Creek China Flat Museum has a Bigfoot exhibit with foot casts, maps, and a 25-foot redwood sculpture.
Northwest Travel, Bigfoot, Sasquatch, Features

redpajama_set_name: RedPajamaCommonCrawl


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Home Page ContentPress Releases SingTel collaborates with Telmap to offer mobile Search, Mapping and Navigation to consumers
SingTel collaborates with Telmap to offer mobile Search, Mapping and Navigation to consumers
by david.nunes September 8, 2010
13 August 2010, Telmap, the global leader in mobile location services, today announced that it has collaborated with SingTel, Asia's leading communications group to offer personal search, mapping and navigation services to its mobile subscribers in Singapore.
SingTel Xplorer, SingTel's first GPS navigation client, enables customers to turn their mobile phones into a powerful navigation system with access to a full suite of rich places of interests (POIs) and real-time traffic updates. Customers can reach their destination efficiently through turn-by-turn navigation with voice or map instructions. Xplorer comes with both Singapore and Malaysia maps.
Traffic data is collected from multiple reliable sources such as the physical locations of SingTel's extensive mobile cell network base and traffic flow feeds. The information is analysed using advanced patented traffic science calculations, which analyses traffic movements of mobile customers and traffic flow. Data is refreshed every five minutes, resulting in high-quality real-time and historical information about traffic flows on most Singapore minor roads.
The application is fully integrated on all GPS supported handsets, including latest mobile operating platforms such as iPhone OS4, Android and Microsoft OS.
"The Asia Pacific market will play an important role in Telmap's global strategy to lead customized and localized Location Based Services across the globe" said Oren Nissim, CEO of Telmap. "We are happy to partner with SingTel in making location services available through their network. Telmap will continue to work closely with mobile operators and support them in offering a competitive local and valuable experience to their subscribers as part of their core services."
SingTel Xplorer offers a rich search, mapping and GPS navigation experience with features including real-time, turn-by-turn in-car navigation with 3D maps and clear voice instructions pronouncing street names and a dedicated guidance mode for pedestrians. SingTel Xplorer also includes live traffic updates which allow users to select their route based on traffic, automatically view traffic events on the map along the route and Xplorer automatically suggests alternative routes while driving based on traffic congestion.
SingTel customers will benefit from Telmap's localized and customized solution, which uses a dynamic content deck, offering dynamic integration of local and premium content to their subscribers, making Xplorer relevant and up-to-date for the local market. SingTel subscribers will benefit from access to relevant local content including places and businesses, restaurant listings, hotels, and more. In the very near future more content will be available such as inSing, global maps, places of interest (POIs) and more.
"Location and navigations services are becoming more popular with customers and it is essential that SingTel provides a reliable and up-to-date solution. SingTel Xplorer will deliver a truly localised experience and will provide relevant traffic information to people on-the-go with our comprehensive traffic data feeds. With various mobile platforms to support Xplorer,
our customers can enjoy a richer, mobile navigation experience," said Ms Ying Lai Chang, Vice President Consumer Products, SingTel.
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An excited client/customer is my main goal and motivation behind every custom website design project I take on. As a professional web designer, nothing is better than hearing that I've accomplished a client's goals, visions and needs just as or better than imagined. Below, I've included testimonials from clients whose projects are also featured and explained in detail in my website portfolio.
" Tony, you're the man!!! Thanks for your magic on my Ohana Athletic Club website!!! Your a GENIUS!
I strongly recommend 3MultiMedia.com Web Design & Hosting to all of my business associates. Tony, the owner, did a great job creating my volleyball camp website. Tony implemented the changes I needed made. He took care of all of my website issues and I can be a very picky customer. Their turnaround time is phenomenal and Tony is someone you can trust and is incredibly accessible via phone. "
" I had a great experience with Tony and 3MultiMedia.com. His professionalism design ability and service level were all far beyond what I expected. He was able to improve on my own vision and bring many ideas to the table. The final product increased my traffic and dropped my bounce rate from around 45% all the way down to 25%. The number and quality of the leads I have received from the new website has all gone up. Within a week of having the new site up I had new clients walking in the door. Not only does the new site look good and function smoothly but it does what counts; get customers. The service, professionalism and results far exceeded my expectations. "
" I had put off getting a website for my business for years, until I finally pulled the trigger with Tony and 3MultiMedia.com. Tony is a true professional and knew exactly where to take my website. Tony had my website online one week after I met with him. The website meets every single one of my needs and looks great. Thank you Tony!!!"
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" I'll definitely send a HIGH recommend your web site expertise to all of the the people that are interested in your services anytime. You did a great web site for www.SPLASHPOOLSNSPAS.com THANKS...BOB DELMONT"
" I came to Tony with an idea and he turned my vision/idea into a great soccer website. Thank you again for everything! Your design professionalism was 1 in a million! I really love my background and how my site flows with Soccer info. - Josh"
" The Edison High School Foundation website is the best! Thank you for taking the time to redesign our site so that we can edit it ourselves! The layout is perfect and not overdone! We have had the pleasure of working with 3MultiMedia (Tony Eden) for the past few years. He has been very helpful in assisting the Edison High School (Huntington Beach CA) Foundation website up our own corner of the web. His approach is always professional and his ideas are constructive and give guidance in our meetings. We would completely recommend Tony @3MultiMedia.com for website design!"
" My Trading Investing site has taken off since Tony at 3MultiMedia did the redesign last year.\ We appreciate your one-to-one work relationship along with your work ethic. You've taken our site to a new level!.It's almost impossible to find professional website designers that are as cooperative, knowledgeable and professional as Tony is."
" Tony and 3Multimedia are great to work with, They listened to my needs and were able to design my perfect website. Not being a techie they explained everything so I was able to understand it. Highly recommend his services and talents. I have worked with several different people and it just seemed like having anyone else do my website would be a pain—luckily I found Tony ~ Don Petry "
" Tony, I just wanted to say thank you for the unbelievably professional design of my web site. I always get this question: "Who did your site?" The way you designed my site is unbelievably done to the most professional degree. You are the best and I know that I can trust you to always work on my updates as well. I would recommend Tony for anyone who is looking for a website."
" Our gaming company's website rocks! Thanks for taking it to the next level. We were all blown away at the design and functionality of the site This is a better website than any other gaming company that I have ever seen..... Thank You again for designing it so well! I would recommend Tony from 3MultiMedia.com any day of the week for any of my friends and or family."
"I could not be more pleased with our business website! Thank you so much for taking my vision and making it happen! Thanks so much for offering advice while still making sure the final product is "mine", I will be sure to recommend you to anyone who asks! Everyone who visits the site gives me nothing but compliments! People are loving my site! You did a superb job. Not only that, but I also love that after just a few weeks in business, I already show up on the first page of a Google search! Tony is a great guy and has a ton of knowledge in website design.


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Home News Sports "I was cleaning gutters for 20 Naira" – Footballer, Victor Osimhen reveals
"I was cleaning gutters for 20 Naira" – Footballer, Victor Osimhen reveals
by Awesome Media October 18, 2021
Super Eagles star, Victor Osimhen has revealed he cleaned gutters and did other menial jobs to survive before venturing into professional football.
The 22-year-old Napoli star shared his grass to grace story while speaking with Head Topics recently.
His words,
"Most people are familiar when I say this because they have come through even worse [situations] than mine and where I grew up, nothing is promised there because whatever you want to get, you have to work so hard for it.
"I remember when I was growing up, I was cleaning the gutters for 20 Naira, for my landlord and did some cleaning jobs for my neighbours close to my house, fetch water for 80 Naira, I think.
"I found pleasure in doing these things because I believed when I worked so hard to get this money, it was important for me. I had consciousness on how to spend this money and help my family.
"While growing up, my brother was selling newspapers, my sister oranges. I was in the traffic with water and I'd had to chase cars, give them water to get the money. I feel like these struggles helped me in life.
"I am grateful for where I am today because of what I'd been through, it shaped me into the man I am. I am really grateful to God for that."
Speaking on his journey to stardom, the striker said, "It all started like a dream, a boyhood dream from the slum and I feel like I have achieved this feat, it's one of my biggest achievements in life."
"I was cleaning gutters for 20 Naira" – Footballer, Victor Osimhen reveals Reviewed by Awesome Media on October 18, 2021 Rating: 5
Tags : News Sports


Pensioners who move abroad need to manage their finances in a manner that will enable them to achieve retirement bliss through their savings and pension income.
However, to successfully become expatriate pensioners, people must first find a suitable location, and make sure they have enough money saved up to enjoy a comfortable retirement.
Another statistic, albeit a more worrying one, is that a large amount of retirees are not seeking out financial advice. Financial advice is incredibly important, especially for expats. There are a number of financial products and pensions schemes that are available to expat retirees, and in some cases you may not be aware of these if you do not speak to an adviser.
For expat retirees, a great option is the QROPS. QROPS, Qualifying Recognised Overseas Pension Scheme, is a pension transfer available to all people who have a UK pension. It is completely recognised by HMRC and is a legitimate way to maximise your existing pension, and thus get more out of your foreign retirement.
If you are planning a retirement abroad and wish to learn more about QROPS, and other options available, get in touch with a recommended IFA.


Brazilian pair ruled out of Sunday warm-up vs Croatia - SABC News - Breaking news, special reports, world, business, sport coverage of all South African current events. Africa's news leader.
Renato Augusto, the Beijing Guoan midfielder who has been a regular starter under coach Tite, will stay in London to have physiotherapy on an injured left knee.
Renato Augusto and Douglas Costa will miss Brazil's World Cup warm-up friendly against Croatia in England this weekend due to injury, the Brazilian Football Confederation said on Friday.
Costa will travel with the squad to Liverpool, but will not play in Sunday's match due to a thigh strain.
However, Renato Augusto, the Beijing Guoan midfielder who has been a regular starter under coach Tite, will stay in London to have physiotherapy on an injured left knee.
Brazil play Croatia at Anfield on Sunday in what will be their penultimate warm-up before heading to Russia.
The five-time World Cup winners also face Austria in Vienna on June 10.
Striker and talisman Neymar, who has not played since injuring his foot in a league match for Paris St Germain back in February, is expected to feature in both games, although reports said he will start on the bench against Croatia.
Brazil kick off their World Cup campaign against Switzerland on June 17 and also play Costa Rica and Serbia in Group E.


Blanket stitch is a wonderful stitch for open and large motifs. You can use the blanket stitch to fill areas completely or partially. The smaller the space between stitches, the more of the underground fabric you will cover. Blanket stitch is a very economic stitch, leaving not much thread in the back, but a lot on the front. It works best with rounded shapes, like ovals or circles, but also scalloped edges.
You can use the blanket stitch for other shapes, too. Here is a version to create leaves. Blanket stitch works great for bigger shapes. Teeny tiny things are not it's strong suit, so go bigger with them!
If you enjoy flower embroidery, you might like my newer flower series with video tutorials! Hop over here to see the full list of flowers.
This post is part of the Flower Embroidery days. Go to this post for the links to all flower embroidery tutorials of this series.