text:
\section{Introduction}

Deformation of crystalline materials depends upon the deformation history undergone by the materials. This is particularly apparent through the existence of hysteresis curves when alternating the loading direction in tension and compression as in cyclic deformation \cite{Bauschinger:1886, Brown:1971, Argon:2008}. A part of the deformation is reversible and the resulting hardening is much smaller that in continuous monotonic deformation, explaining why deformation can be easily accumulated this way. At every two cycles, the flow stress may increase usually in a continuous manner, until reaching a \emph{saturation shear stress} \cite{Cheng:1981, Mughrabi:1978,Lepisto:1986,Bretschneider:1997, Gong:1997,Li:2009, Li:2011}. In single crystals, this saturation stress depends upon the loading direction and the materials as it not always scales with the shear modulus \cite{Li:2011}. For a given loading direction and material, the saturation stress typically increases with the cycle shear increment $\gamma_{p,cy}$, and the saturation stress may well be below the stress values observed in monotonous deformation. For example, in Cu -certainly the most studied system- deformed in single glide condition, the cyclic saturation stress ranges from 16 to 40 MPa exhibits a three stage curve with a central plateau value of 28 MPa. In stable multislip conditions, the cyclic saturation stress typically increases monotonously with an apparent slope that depends upon the loading direction and the material \cite{Bretschneider:1997, Gong:1997,Li:2009, Li:2011} and may reach 50 or 60 MPa. In comparison the maximum stress reached by monotonically deformed single crystal of Cu is in the range of {80-100} MPa.

Until recently, this tension-compression asymmetry -also known as the Bauschinger Effect (BE)- was commonly thought to be related to the building up of Long-Range Internal Stresses (LRIS) or backstress associated to the formation of dislocation patterns according to the so-called composite model \cite{Asaro:1975eu, Mughrabi:88, Argon:2008}. However, dislocation patterns are rather weak at smaller strain, they may well be different from the archetype of the composite model, and no LRIS was found in recent large scale mesoscale simulations \cite{Queyreau:2021} or by X-ray microdiffraction \cite{Kassner:2009fk, Kassner:2013fk} (except at very large strain \cite{Levine:2019}). Another explanation consisted in the partial dissolution of the microstructure \cite{Buckley:1956, Sleeswyk:1978, Depres:2008fk, Rauch:2011kx}, as a mechanistic way to reduce dislocation density and thus flow stress. However, no clear dislocation elementary mechanism was clearly identified as there is no in-situ observation of dislocation motion during cyclic deformation. Ultimately, these explanations explain neither the transient nature of the Bauschinger effect nor its reversibility component.

The present authors \cite{Queyreau:2020,Queyreau:2021} recently proposed a systematic study of cyclic deformations in single crystals by means of Discrete Dislocations Dynamic (DDD) simulations. Despite of the present of a pronounced Bauschinger effect in the simulations, no LRIS was measured in the simulations. Statistical analysis of the DDD simulations showed that the tension-compression asymmetry is caused by two original elementary mechanisms of dislocations. The first mechanism i) is related to easy destruction of binary junctions as their stability is asymmetrical as they formed from mobile segments whose curvature is driven by the applied loading. Junctions formed is tension are thus more stable in tension than in compression, and are easily destroyed during the backward motion of dislocation in compression. The second elementary mechanism ii) is related to the reduction of the mean free path of dislocations as they glide in regions of the crystal that they already explored. In the backward motion, the mobile segments unwind stored segments on the edge of the swept area, leading to a reduction in the storage rate. These mechanisms naturally explain the transient nature of the tension-compression asymmetry and the reversible component of plastic deformation. From these results, the authors proposed a modified Crystal Plasticity framework, that was implemented into a FEM to include the physics highlighted in the DDD. As a results this multiscale approach was successful in predicting the Bauschinger effect, hysteresis and saturation stress reported in most of the experimental literature on cyclically deformed fcc single crystals.

Now that we have a CP framework justified from mesoscale physics and validated by confrontation with the existing single-crystal literature, we can now assess the full implications of these results and of the model. In particular, when comparing the literature data obtained though out decades on cyclic deformation of single crystals - the saturation stresses - observed in cyclic deformation are clearly more reproducible that the Bauschinger effect observed on a single tension-compression experiment. This fact suggests that the saturation stress in cyclic deformation is the sole result of the average of some basic dislocation mechanisms, while the Bauschinger effect is impacted  by the initial state of the material, such as the impurity content, the initial dislocation density or the microstructure, but also by the loading conditions. In other words, the saturation stress seems to be a basic property of the material and dislocations. Simulation results could also be processed in order to provide simple equations to help in interpreting and/or fitting the experimental data or to trace back to some of the fundamental constants of the CP model e.g. in an inverse approach. Finally, the single crystal is an integral part of the more complex polycrystal system. The elementary mechanisms described above are general enough to be operative in a polycrystal. The insertion of GB adds a new lengthscale, and a competition between this lengthscale and the ones associated with elementary mechanisms i and ii. These ideas will act as motivations of the present paper.

The objectives of the present work are thus to extend and analyse the CP FEM results obtained for the cyclic deformation of fcc single-crystals. We will see that the saturation stress of cyclic loading is actually related to the theoretical saturation stress obtained in monotonous conditions. Analytical solutions to the constituting differential equations of the CP framework will be provided and can be employed to help in interpreting experimental data. Finally, the implication of the mechanisms i and ii on the cyclic deformation of polycrystal will be assessed.

\section {Methodology}

This section presents the Crystal Plasticity framework derived from the DDD results in \cite{Queyreau:2021}. Readers interested in the presentation of the technical details of the DDD technique along with a review of recent progresses are referred to \cite{Devincre:2011fk,Queyreau:BDM}. This CP model is an extension of the dislocation density based  descriptions proposed in the seminal work \cite{Kubin:2008fk, Kocks:03} and subsequent works inspired from DDD \cite{Devincre:2006uc} for monotonous loadings. In fcc metals at intermediate and high temperature, the flow stress is controlled by the formation and destruction of reactions -junctions- between dislocations belong to different slip systems. The critical shear stress $\tau_c^i$ on the active slip system 'i' is thus related to the \emph{forest} obstacle densities $\rho^j$ and expressed as :
\begin{equation}
\tau_c^i = \mu b \sqrt{ \sum_j a_{ij} \rho^j}
\label{eq:forest}
\end{equation}
Where $\mu$ is the isotropic shear modulus, $b$ the norm of the Burgers vector and $a_{ij}$ the components of the interaction matrix. These last interaction coefficients measure the average strength of the two interacting slip systems and can be determined in a straigthforward manner from DDD \cite{Devincre:2006uc, Queyreau:2009lr,queyreau:2010, Madec:2017kx}. For reasons of symmetry, only six different interactions of different nature exist among the twelve $a_0/2 <110>{111}$ slip systems existing in fcc metals. The interactions are the self-interaction, the coplanar interaction, three junctions: Lomer, Hirth and glissile reactions, and the collinear interaction. The interaction matrix has to have non identical components in order to reproduce fully the anisotropy of the plastic deformation of fcc single-crystals.

A first extension of the CP framework  was introduced to capture the asymmetry of dislocation junctions (mechanism 'i' in the introduction) and to reproduce the tension-compression asymmetry observed in DDD \cite{Queyreau:2021}. Junctions formed during the forward loading from curved segments inherit a mechanical stability asymmetry. The junctions formed during prestrain or forward tension loading, are stronger when stressed in tension and statistically weaker when now stressed in compression, and this asymmetry in mechanical stability can be considered as a line tension effect, with the parent moving segments colliding when they were in extension. As a results, while plastic flow is still controlled by junctions destruction, the interaction coefficients during the backward loading are effectively weakened over a transient that depends on the amount of the initial forward deformation. The pool of weak junctions gets destroyed as  backward deformation proceeds, until mobile segments explore new area of the crystal and form junctions polarized in compression. A reversibility function $r_a$ was introduced to reproduce these effects as:

\begin{equation}
a_{i j}^{b c k} = ( 1 - r_a ) \times a_{i j} \mathrm{, ~with ~~} r_a = \exp \left(- \frac{ \gamma_{b c k}^i} {C_a b \sqrt{ \Delta \rho_{pr}^i}} \right)
\label{eq:newa}
\end{equation}

with $a_{i j}$ the reference and constant interaction coefficient measured in continued tension. the subscripts 'pr' and 'bck' refer to prestrain and backward loading, respectively. In a cyclic loading, the prestrain  becomes the previous cycle, while the backward loading corresponds to the current cycle. The amount of backward strain impacted by this transient on the junction stability is approximated as $C_a b \sqrt{\Delta \rho_{pr}^i} $. This last equation can provide a lengthscale estimate to this effect. The density increase during previous deformation cycle is  ${ \Delta \rho_{pr}^i} = \max(\rho_{pr}^i) - \min (\rho_{pr}^i)$ and corresponds to the potentially impacted dislocation density polarized according to the initial loading. The constant $C_a$ has been determined as $C_a = 0.6 \pm 0.1$ through a statistical analysis of the interaction coefficients over a large panel of relevant DDD simulations \cite{Queyreau:2021}. The $\gamma_{b c k}^i / C_a b \sqrt{ \Delta \rho_{pr}^i}$  term within the exponential of $r_a$ states the competition between the easy destruction of junctions formed in tension, and the formation of new junctions in never-explored regions of the crystal.

A second fundamental equation in the CP framework describes the evolution of the dislocation density on active slip system 'i' with the system shear $\gamma^i$. The evolution of the dislocation density is related to the kinetics of plastic activity that occurs through intermittent busts or avalanches of dislocation motion. In principle, dislocations are stored at the end of an avalanche as dislocation segments left at the edge of the swept area. This being said the exact theoretical connection between dislocation avalanches at the mesoscale and the observed continuous macroscale storage of dislocations has still to be formulated. For the present work and in monotonic loading conditions, the density evolution takes the following simple form:

\begin{equation}
\frac{d \rho^i}{d \gamma^i} = \frac{1}{b} \left(  {1 \over L_{hkl}}  + {1 \over L_I} - y \rho^i \right) = \frac{1}{b} \left(  \frac{\sqrt{\sum_{i \neq j} a_{ij} \rho^j}}{K_{hkl}}  + \frac{\sqrt{a'_0 \rho^i}}{K_I} - y \rho^i \right)
\label{eq:KM}
\end{equation}
where the first two terms relate to the dislocation storage and the last term represents the dynamic recovery. The storage rate is commonly related to the Mean Free Path (MFP) of dislocations $L^i$, which represents the average distance covered by dislocations before the temporary or permanent storage. The MFP of dislocations typically depends upon the loading axis and thus the number of active slip systems. In \cite{Devincre:2006uc,Kubin:2008fk}, Kubin et al. proposed a simple decomposition of the MFP into elementary ingredients: i) the rate $p_0$ of forming a junction that will ultimately store a dislocation segment, and ii) the average length of stored dislocation $<l>$ and iii) junction segments and $<l_j>$, both of which scale with the dislocation density $<l> = k_0 / \sqrt{\bar{a} \rho}$ and  $<l_j> = \kappa <l> $. $\bar{a}$ is the average interaction coefficient for the considered orientation. Statistical analysis of DDD results showed that parameters $k_0, p_0, \kappa$ to be constants independent of the loading direction. The MFP can thus be written:
\begin{equation}
L^i = \frac{\mu b^2}{\tau_c^i} \left[ \frac{\sqrt{\bar{a}} n (1 + \kappa)^{3/2}}{p_0 k_0 ( n - 1 - \kappa )} \right]
\label{eq:MFP}
\end{equation}
the third term is associated to the dynamic recovery occurring at large dislocation density. $y$ is related to critical distance at which two dislocations can easily annihilate, which may be measured from atomistic data. Here, $y$ takes different values depending on the loading axis to reproduce the anisotropy of the onset of stage III observed on experimental deformation curves on single-crystal \cite{Takeuchi:1975, Kubin:2013fk}.

\begin{table*}[!ht]
\caption{List of physical parameters employed in the CP simulations of the cyclic deformation of fcc single crystals. Most of these parameters are coming from DDD simulations for monotonic or cyclic deformation, few parameters are coming from the experiemental litterature but for monotonic loading only. No parameters or equations were fitted to the cyclic deformation literature. The resulting FEM simulations are therefore truly predictive of cyclic deformation. }
\label{tab:paramCP}

\begin{center}
\small
\begin{tabular}{ c c c c c c c }
\hline
 ${a'_0}$ (self) &   ${a_{ortho}}$ (Hirth)  & ${a_2}$ (glissile) & ${a_3}$ (Lomer) & ${a_{colli}}$ (Collinear) & ${\rho_{ref}}$ & ${\rho_{0}}$\\
 0.122 & 0.07 & 0.137 & 0.122 & 0.625  & $10^{12}\,$ m$^{-2}$ & $10^{12}/n_{sys}$  \\
 \hline
  ${K_{I}}$ & ${K_{112}}$ & ${K_{111}}$ & ${K_{001}}$ & ${C_a}$ & ${A_p}$ & ${C_p}$ \\
  90 & 10.42 & 7.29 & 5 & 0.6 $\pm$ 0.1 & 2 $\pm$ 0.6 & {2.3 $\pm$ 0.3} \\
 \hline
 ${\dot{\gamma}_{app}}$ & ${\dot{\gamma}^i_{0}}$ & m &  ${y_I}$ (SG, [112]) & ${y_{001}}$ (nm) & ${y_{111}}$ (nm) \\
 $n_{sys} \times {\dot{\gamma}^i_{0}}$ & 10$^{-4}$ s$^{-1}$ & 35 & 0.5 nm &  {3.6 (Ni), 3.4 (Cu)} & {2 (Ni), 1.5 (Cu) }\\
 \hline
\end{tabular}
\end{center}
\end{table*}

When considering the cyclic deformation, DDD simulations \cite{Queyreau:2020, Queyreau:2021} showed that dislocation evolution is still associated to the storage of segments in the wake of avalanches at the fringe of the area swept by dislocations, with some similarities to what was observed in alloys \cite{Queyreau:09}. Equation \ref{eq:KM} is thus still valid. However, the reduced storage (mechanism ii in the introduction) observed at the loading reversal where the stored segments are simply unwinded, needs to be taken into account through a modified MFP (through a change of rate of locking dislocation rate $p_0$). A second reversibility function $r_p$ is thus introduced \cite{Queyreau:2020, Queyreau:2021} as:
\begin{equation}
p_{0}^{b c k} = ( 1 - r_p ) \times p_{0} \mathrm{, ~with ~~} r_p = A_p \times \exp \left( - \frac{ \gamma_{b c k}^i} {C_p b \sqrt{ \Delta \rho_{pr}^i}} \right)
\label{eq:P0bck}
\end{equation}
Where $A_p$ and $C_p$ are two additional constants measuring the initial MFP drop and the transient length of the reduced MFP, respectively. Statistical analysis over the transient observed in DD showed that $A_a = 2 \pm 0.6$ and $C_p = 2.3 \pm 0.3$. Interestingly, the transient on the reduced MFP (ii) is much larger that the one on the reduced junction stability (i). These different transients explain the non-monotonic response at the beginning of the backward deformation that is sometimes observed on experimental deformation curves \cite{Ebener:1991, Daniel:1971, Nasu:1984}.

Finally, the flow rule provides a close form to the CP framework in connecting the plastic activity $\dot{\gamma^i}$ on the active system to the critical resolved shear stress $\tau_c^i$  and the applied shear stress $\tau^i$:
\begin{equation}
\dot{\gamma^{i}} = \dot{\gamma_0^{i}} \left(\frac{\tau^{i}}{\tau_c^{i}} \right)^{m}
\label{eq:strainrate}
\end{equation}
In the case of fcc metals at room temperature, the  strain rate sensitivity is related to the formation and dragging of jogs along dislocation lines. Here, the constant $\gamma_0^i$ is taken as $\gamma_{app} / n_{act}$, with $\gamma_{app}$ the applied shear rate and $n_{act}$ the number of active slip system. The sensitivity exponent is chosen as $m > 35$ to stay as close as possible to the Schmid criterion. The proposed CP modeling focuses on the some fundamental mechanisms of dislocations that are described by some average constants obtained from statistical analysis of DDD results. The resulting FEM simulations are expected to be general and representative of fcc pure systems. The simulations  will not include plastic localization nor finite geometry effects for now.

Great care has been paid to the numerical resolution of the set of ODE presented above using the Z-Set Finite Element software and Matlab. We employ a double nested Newton-Raphson implicit scheme to solve the corresponding non-linear equations and obtain $\rho^i$ and $\tau^i$ for a given time step. Gradients required in the Newton-Raphson are expressed analytically, and a relative convergence criterion was set to $10^{-7}$. The representative volume was set to a single point of integration in order to simulate several thousand of cycles of deformation and finite geometry effects are left for a future work. The parametrization and resolution strategy were validated in \cite{Queyreau:2021} by the one-to-one comparison of CP predictions with the reference DDD simulations.

\section{Results}

\subsection{Saturation stress in monotonic deformation in absence of damage}

\begin{figure}
    \centering
    \includegraphics[width=1\textwidth]{Figure1_abc.pdf}
    \caption{a. One-to-one comparison between the CP prediction (thick continuous lines) and the reference data from Takeuchi \cite{Takeuchi:1975} (circles) on the monotonous deformation of Cu single-crystals at RT. Key single-crystal orientations are chosen. b. Extension of the CP prediction until true saturation of the stress is obtained, in the absence of finite geometry and plastic localization effects. c. Corresponding total dislocation density evolutions in log-log scale. Note that the saturation density for [001] and [111] simulations are of the order of $> 10^{15}$ m$^{-2}$, which is still lower than the saturation density obtained in recent large scale MD simulations of single-crystals of Al \cite{zepeda:2021}.
    }
    \label{fig:fig1}
\end{figure}

In this first section, we focus on the basic case of the single-crystal deformed monotonically in tension. This constitutes as well-pose problem, for which reference experimental data exists. The plastic response is mostly homogeneous during most of the deformation, in the sense that the slip activity is similar in all part of the single-crystalline sample \cite{Kahloun:2016}. We will see later that this basic plastic response can be connected to cyclic deformation.

First, to demonstrate the validity of the CP framework and its parametrization from DDD results, we start with a one-to-one confrontation of the CP prediction with some of the reference data from the experimental literature on single crystal deformation \cite{Takeuchi:1975}. Such experiments are rather delicate to perform as the materials state and experimental conditions can have a dramatic impact on the mechanical response, such as the impurity content, the precise orientation of the sample, or whether the jaws can rotate to accommodate crystalline rotation and ensure uniaxial deformation. Here, we compare our CP simulations with the reference work from Takeuchi on Cu single-crystals for selected loading directions: [112], [111] and [001] directions. The initial density was slightly adjusted in the simulations about $10^{12}$ m$^{-2}$ to capture the initial flow stress.

This comparison is shown in Figure 1.a, where a nice quantitative agreement is found  between the simulations and the experimental reference data. The classical picture of the response of the single crystal is recovered here. The initial hardening rate depends upon the orientation and the number of slip systems activated simultaneously, e.g. in a increasing order: [112], [111] and [001] having two, three and four slip systems activated simultaneously. For [111] and [001] loading directions, these numbers of slips correspond to only half of the six or eight possible slip systems as these are pairs of colinear interactions. The colinear interaction is known to be a specific reaction among junctions, as it leads to the annihilation of the intersecting segments of dislocations and is associated to a very large hardening due to the shortening of mobile segments. DDD simulations have shown that when starting with pairs of colinear systems, one of the colinear system will take over the other for [111] or [001] simulations \cite{Devincre:2007}. Another notable aspect of the single-crystal response is the anisotropic onset of the stage III, when the dynamic recovery becomes noticeable. The dynamic recovery of the [111] in particular is weaker, so much so that the [111] curve eventually crosses the initially steepest [001] curve at a deformation of about 40\% of plastic strain.

In agreement with the response of ductile materials, prior to the fracture of the sample, the experimental flow stress decreases after a maximum corresponding to the striction of the sample. This decrease is expectedly absent in the CP simulations as plastic localization and fracture mechanics are not included. Fracture mechanics may be included in different manner in FEM, for example using a cohesive zone framework \cite{Mishin:2010,Rice:1989}, in connection to atomistic mechanisms controlling the debonding of matter \cite{Vanderven:2004, Ehlers:2016,Ehlers:2017}. For now, let us focus on the plastic response in the absence of fracture mechanism. For a deformation corresponding to the experimental fracture, the hardening rate in the simulations has still a non-zero value. When recalling that  equation \ref{eq:KM} was proposed initially to saturate, the CP deformation curves are thus expected to saturate at a larger deformation.

We thus expanded the deformation range until saturation of the stress is obtained, and this is shown in figure 1.b for Cu. All curves saturate eventually, the corresponding deformation at saturation is however rather large, even beyond 100\% for the [001] and [111] curves. The saturation stress in monotonous condition  $\tau_{sat,mo}$ follows the following hierarchy, with the single glide condition [135] having the smallest saturation stress of 28 MPa, [112], [001] and [111] having the largest $\tau_{sat,mo}$ = 140 MPa. Since $y_c$ is material independent for the [135] and [112] curves, the $\tau_{sat,mo}$ for these curves scales nicely with the shear modulus of the considered materials. This saturation stress values will be exploited a bit further.

From a theoretical point of view, this saturation stress observed in the simulation can be predicted from the set of constituting ODE of the CP framework. The starting point is obviously the dislocation density evolution that has to be set to zero. In the case of single glide the resolution is straigthforward leading to:
\begin{equation}
    \rho_{sat,I} = {\sqrt{a_0 '} \over K_I^2 y_{c,I}^2}
\end{equation}
The stable multislip conditions are a bit less straightforward to solve as the interaction coefficient exhibit a dependence upon the dislocation density inherited from the simplified line tension that is assumed in the critical stress equation \cite{Devincre:2008lr,queyreau:2010}. When the dislocation density is large this correction becomes important, which is the case at the large saturation stresses under consideration here. The logarithmic correction is defined as $c(\rho^i) = \ln \left( 1/ b\sqrt{a_{ij} \rho^j} \right) / \ln \left( 1/ b\sqrt{a_{ij} \rho_{ref}} \right)$, where $\rho_{ref}$ is the reference dislocation density used to determine the interaction coefficient $a_{ij}$. The saturation density is now:
\begin{equation}
    \rho_{sat,hkl} = \left( {\sqrt{a_0 '} \over K_I y_{hkl}} +  {\sqrt{ (n_{act}-1) \bar{a} c(\rho^i)  } \over K_{hkl} y_{hkl}}  \right)^2
\end{equation}
where $n_{act}$ is the number of activated systems. This equation is an implicit equation for $\rho^i$, but since the density is contained in a logarithm function, this equation converges after only few iterations to evaluate $\rho_{sat}$.  Finally, the saturation stress is obtained in inserting the $\rho_{sat}$ into the critical stress:
\begin{equation}
    \tau_{sat} = \mu b \sqrt{\bar{a} n_{act}\rho_{sat}}
\end{equation}
In the past equations, the saturation stresses are solely function of interaction coefficients, MFP and $y$. In the model, these quantities are defined from statistical averages over evolving dislocation microstructures, and are weakly impacted by the nature of the fcc system under consideration (at least not as a  first order approximation).  The initial dislocation microstructure or density are absent from these equations. The saturation stresses appear thus as a \emph{fundamental dislocation property} representing the balance of storage and recovery processes among dislocations in fcc crystals.

An approximate solution for the evolution $\rho^i (\gamma^i)$ is provided in the appendixes for single and stable multislip conditions. Theses functions can be used to interpret or fit experimental data. In monotonous condition, the dislocation density increases monotonically until saturation. The strain corresponding to the saturation density can be estimated from the density evolution as an integral:
\begin{equation}
\Delta \gamma^i = \int_{\rho_0} ^{\rho_{sat}} \left[ \frac{1}{b} \left(  \frac{\sqrt{\sum a_{ij} \rho^j}}{K_{hkl}}  + \frac{\sqrt{a'_0 \rho^i}}{K_I} - y \rho^i \right) \right]^{-1} d\rho^i
\end{equation}
This only provides an estimation as the saturation density is only reached as $\gamma^i$ tends to $+\infty$. We now recover a dependence upon the initial dislocation density $\rho_0$ that affects the amount of deformation required before saturation.

\subsection{Saturation stress in cyclic deformation in single crystals}

\begin{figure}
    \centering
    \includegraphics[width=1\textwidth]{Figure2_abc.pdf}
    \caption{
    The saturation shear stress $\tau_{sat,cy}$ for cyclic deformation in fcc metals measured in a number of experimental studies \cite{Cheng:1981, Mughrabi:1978,Lepisto:1986,Gong:1997,Bretschneider:1997, Li:2009, Li:2011} compared to the prediction of the CP model in \cite{Queyreau:2020,Queyreau:2021} and to the theoretical saturation stress in monotonous conditions from previous section. Data is shown as function of the loading direction: a) initially single glide condition, b) [001] and c) [111] for Cu, Ni and Ag and the plastic strain increment $\gamma_{p,cy}$
    }
    \label{fig:fig2}
\end{figure}

The CP framework can be applied to alternating loadings and the simulations capture most of the features of cyclic deformation \cite{mughrabi:2010,Li:2011}. The maximum stress reached at each cycle increases monotonically until spontaneously saturating. For example in single glide conditions, we recover the so-called three stage curves of the now saturation stress in cyclic deformation $\tau_{sat,cy}$ with respect to the shear increment per cycle $\gamma_{p,cy}$. This is shown in figure 2.a for some fcc metals in comparison with most of the existing experimental literature \cite{Cheng:1981, Mughrabi:1978,Lepisto:1986,Gong:1997,Bretschneider:1997, Li:2009, Li:2011}. In the central plateau (or pseudo plateau) the microstructure typically transform from no Persistent Slip Band (PSB) to progressively fully constituted of PSB \cite{mughrabi:2010,Li:2011}. The plateau ends for strain increments of the order of 1\% when a secondary slip system activates. In the CP simulations, the center region of the curve is not exactly a plateau as it exhibits a slight slope.

Figure 2.b and 2.c show the saturation stress for cyclic deformation of Cu and Ni single crystals oriented along [001] and [111] in stable multislip conditions. The saturation stress $\tau_{sat,cy}$ increases here monotonically without the presence of stages. The slope of the experimental curves seems however to decrease as the strain increment per cycle $\gamma_{p,cy}$ increases. This effect is more apparent on the data concerning Ni single crystals. The CP predictions of the $\tau_{sat,cy}$ are in nice qualitative and quantitative agreement with the existing experimental data. The CP predictions seem however to overestimate the saturation stress at larger  $\gamma_{p,cy}$, and this was suggested to be a consequence of the absence of dislocation microstructure in the CP simulations \cite{Queyreau:2021}.

The saturation stress $\tau_{sat,cy}$ can be also obtained from the analytical resolution of the set of ODE presented in section 2. The resolution is however a bit more delicate with the presence of the two reversible functions $r_a$ and $r_p$ that are themselves function of $\gamma^i$ and the  dislocation density $\Delta \rho^i$ stored on the previous cycle. In cyclic deformation, we highlight that the saturation stress does not mecessary correspond to an horizontal tangent of the shear strain-shear stress curves of the considered cycle (except in the large $\gamma_{p,cy}$ limit discussed a bit further). At the end of a cycle at saturation, the hardening may well be non zero, but wince the cycle starts with a large dislocation density decrease, the increment over the entire cycle is null. To obtain the saturation density, one has to solve the integrated density evolution function:
\begin{equation}
{\Delta \rho^i}=0= \int_{0}^{\gamma_{p,cy}/n_{act}} \frac{1}{b} \left(  \frac{\sqrt{\sum a_{ij} \rho^j}}{K_{hkl}}  + \frac{\sqrt{a'_0 \rho^i}}{K_I} - y \rho^i \right) {d \gamma^i}
\end{equation}
Therefore, the dislocation evolution $\rho^i(\gamma^i)$ must be solved first contrary to the monotonous case. This can be done if the logarithm correction is approximated as a Taylor series where the dislocation $\rho^i$ in the current cycle is in the proximity of $\rho^i_{sat}$. The details are given in the Appendixes. The resolution leads once more to an implicit equation to obtain $\rho^i_{sat}$.

Next, we connect the saturation stresses obtained in monotonous and cyclic deformation. Previous section has shown that the saturation stress in monotonous conditions can be obtained for rather large deformation amount. These saturation stress results could thus be drawn as well on the figure 2 for one cycle of the corresponding amount $\gamma_{p,cy}$. One would expect that the $\tau_{sat,mono}$ would act as an asymptotic limit to the $\tau_{sat,cy}$ curves. This is exactly what is obtained when considering the case of single glide condition. The saturation stresses obtained in single glide and monotonous conditions in previous section are strikingly close to the saturation stress at the plateau in cyclic conditions. The $\tau_{sat,cy}$ curves obtained from experiments or from the CP simulations asymptotically converge towards $\tau_{sat,mono}$.

When now considering the third stage of the curves in figure 2.a or the multislip conditions in figure 2.b and 2.c, the agreement with experiments is a bit less obvious. The CP simulations are carried out past the range of $\gamma_{p,cy}$ considered in the experiments. Similarly to the single glide condition, the simulated $\tau_{sat,cy}$ curves converges towards the $\tau_{sat,mono}$ limit defined in previous section. In figure 2.a, the third stage converges certainly toward the $\tau_{sat,mono}$ in secondary double glide, that is probably out of reach. The agreement with the experimental data is rather good for the case of Cu. However, the quantitative agreement between with the model is a bit less good for Ni single crystal data, nonetheless the experimental data seem to enter a plateau for the largest $\gamma_{p,cy}$ considered experimentally, which could be well agree with the analysis proposed here, but not the quantitative values of the saturation stress. The $\tau_{sat,mono}$ being here overestimated in the model due maybe because of some specific parameters less well documented for Ni, such as $y_{hkl}$ or the stronger impact of microstructures that are absent in the model.

\subsection{Saturation stress in polycrystals}

\begin{figure}
    \centering
    \includegraphics[width=0.5\textwidth]{Fig_sigma_sat_polyx.pdf}
    \caption{
    compilation of saturation stresses obtained experimentally on polycrystalline Cu with large grain size \cite{Magnin:1984} as function of strain increment $\Delta \epsilon_{p,cy}$. The data covers the work from Saxena and Antolovich (1975, squares), Figueroa et al. (1981, circles), Mughrabi and Wang (1981, triangles) and Rasmussen and Pederson (1980, diamonds)
    }
    \label{fig:fig3}
\end{figure}

This last section aims at connecting our results on the single crystal to the more complex and general polycrystalline system. What comes next is simply an illustration as the plastic response of the polycrystals is the average of many different features or elementary mechanisms that are not addressed here, ranging from the crystallographic texture, grain size and morphology, and dislocation-grain boundary interactions. Besides, simulating using CP FEM the cyclic deformation over more than 10,000 cycles while accounting for the a representative polycrystalline microstructure seems computationally out of reach, still to date. This being said we reprise a part of the analysis initiated by Magnin et al. in \cite{Magnin:1984}, where they compiled the cyclic response of polycrystalline Cu with relatively large grains. The collection of saturation stresses $\sigma_{sat}$ from the literature is displayed in figure 3. These curves typically show an increase in the $\sigma_{sat}$ with the plastic increment per cycle $\epsilon_{p,cy}$, some curves may suggest a quasi plateau regime starting for $\epsilon_{p,cy} > 10^{-5}$ and ending around  $\epsilon_{p,cy} \approx 10^{-3}$. The deformation microstructures transform from veines to PSB that are typical of the patterns observed in single glide conditions in single crystals. For $\epsilon_{p,cy} >\approx 10^{-3}$, the saturation stress increases more rapidly. The deformation microstructures now correspond to mazes, that are typical of the patterns observed in single-crystals deformed in multislip conditions.

The polycrystalline plastic response can be understood in terms of effective single-crystal response through the so-called Taylor coefficient $M$, which depends upon the crystallographic texture. The so-called Hall-Petch effect is neglected and this is a valid assumption in the case of large grain polycrystals. In the absence of precise $M$ measures, the Taylor or Sachs hypothesises are commonly considered as bounding limits. The figure 4 shows theses two transformations of the polycrystal saturation stress into a single crystal approximation as $\tau_{sat} = \sigma_{sat} /M$. The Sachs is typically a good approximation at small strain, while the Taylor approximation works well at large strain. The average single crystal behaviour is expected to follow Sacks approximation first and transition to the Taylor ones at larger deformation. These curves will be compared with the saturation stress obtained experimentally on single crystals. At small strain, one might expect that grains in the polycrystal deform in single glide condition as suggested by the dislocation microstructures, and we see a sticking agreement between the single crystal $\tau_{sat}$ and the Sachs approximation for $\epsilon_{p,cy} < 5.10^{-4}$. For very small strain, so grains may remain in the elastic regime. For large deformation, one might expect that all grains to deform in multislip conditions, in accordance with the maze microstructure observed experimentally, and the [001] curve in particular remains between the two approximations. Between these two extreme situations, the polycrystalline behaviour is certainly a composite response of the single glide and multislip grain activity, with more and more grains transitioning into multislip as $\gamma_{p,cy}$ is increased.

\begin{figure}
    \centering
    \includegraphics[width=0.5\textwidth]{fig_tau_sat_polyx.pdf}
    \caption{
    Comparison of the cyclic saturation shear stress of single crystals and polycrystals in pure Cu from Mughrabi \cite{Mughrabi:1981} and Wang  using the Taylor coefficient following the procedure from Magnin et al. \cite{Magnin:1984}.
    }
    \label{fig:fig4}
\end{figure}

These correlations may thus explain the shape of the raw $\sigma_{sat}$ curves from the previous figures. Most grains of the polycrystals are thus following the single crystal response, which will eventually reach a saturation stress corresponding to $\tau_{sat,mono}$ of 28 MPa. The plateau in the single-crystal system is transformed in a quasi-plateau in the polycrystalline system as more grains transition into multislip conditions.  At large $\epsilon_{p,cy} =\epsilon_{III}  \approx 10^{-3}$ the polycrystal enters stage III where most of the grains are deformed in multislip conditions. Interestingly, the curves for single slip and [001] obtained for the single crystal cross for a deformation $ \approx M \gamma_{III} = M*3e{-3}$ that corresponds well to the onset of the third stage $\epsilon_{III}$ observed in the figure 4. The cyclic behaviour of the polycrystal with large grains seems thus to correspond simply to the average of the individual single grain behaviour, similar to what is accepted for polycrystals in monotonic deformations.

Finally going back to our CP framework, we showed that its predictions were in quantitative agreement with the single crystal cyclic response, and we proposed expressions to predict the saturation stress in these conditions. The previous qualitative analysis means that we can use these results to estimate the saturation stress of the polycrystal with large grains as well.

\section{Conclusion}

In this paper, we employ a physically based CP FE model derived from our recent DDD analysis \cite{Queyreau:2020,Queyreau:2021} to analyse the saturation stresses obtained in monotonic and cyclic deformations of single-crystals and polycrystals of fcc metals.

\begin{itemize}
    \item First, we compare the model results to reference strain-stress deformation curves from the experimental literature on Cu single-crystals. In the absence of plastic localization and fracture mechanism, the CP model expectedly saturates at large strain.
    \item For cyclic deformations, the CP model reproduces the cyclic saturation over various loading conditions and fcc metals. The saturation stress observed in monotonic conditions acts as an asymptotic behaviour at large cycle increment strain $\gamma_{p,cy}$ to the cyclic deformation. This asymptotic behaviour is clearly reached in the plateau stage of single-glide conditions for several fcc metals and for Ni single crystals in multi-slip conditions at large strain increment.
    \item From an analysis of the experimental literature on large grain polycrystals, we have shown that the polycrystalline response to cyclic deformation corresponds to the response of an aggregate of effective single-crystals. At small strain increment, grains are deformed elastically or in single slip, a quasi plateau can thus be sometimes seen as in single-crystals. Then at larger strain, grains transition into multislip conditions with larger saturation stresses.
    \item We proposed analytical or approximated solutions of the CP ODE to predict the saturation stresses in single or multislip conditions for single crystals. These models can be employed to interpret experimental data.
    \item These results were obtained in the absence of plastic localisation and dislocation microstructure, which means that these features have only a secondary order impact on the macroscopic behaviour.
    \item The saturation stress in monotonic or cyclic conditions appear as a fundamental property of dislocations mechanism as it is simply related to averages of dislocations interactions (through the interaction coefficients), dynamical recovery mechanism (through $y$), and reversibility part of the microstructure.
    \item These quantities are phenomenological averages from DDD, where large number of binary interactions are occurring in absence of patterns.
    \item Experimental data on the saturation stress can thus now be used to define these physical quantities and the procedure can be applied to other bcc or hcp crystalline systems.
\end{itemize}

redpajama_set_name: RedPajamaArXiv


text:
Q: Pass Method as parameter for another Method in Java I am creating a generic method to create action listeners for buttons, and i would like to be able to pass in another method into it's actionPerformed, rather than creating a new method for each use case of a listener. something like this:
public static ActionListener makeListener(Method method) {
    ActionListener listener = new ActionListener() {

        @Override
        public void actionPerformed(ActionEvent e) {
            method();
        }
    };
    return listener;
}

Other similar posts have had recommended Lambda Expressions, but (maybe due to a lack of understanding), I don't think those are what i need. Are there other options?

A: Sounds like a Runnable is exactly what you need.
Declaration:
public static ActionListener make_listener(JButton button, Runnable action) {
    ActionListener listener = new ActionListener() {

        @Override
        public void actionPerformed(ActionEvent e) {
            action.run();
        }
    };
    return listener;
}

Or using a lambda for the ActionListener (and skipping the local variable):
public static ActionListener make_listener(JButton button, Runnable action) {
    return e -> action.run();
}

You can now call it with any zero-argument method using a lambda or method reference:
ActionListener listener1 = make_listener(() -> callMethod());
ActionListener listener2 = make_listener(this::callMethod);

redpajama_set_name: RedPajamaStackExchange


Is Nintendo the only one coming with really unique gaming soundtracks?
Forums - Gaming Discussion - Is Nintendo the only one coming with really unique gaming soundtracks?
Who makes the best gaming soundtracks these days?
And I'm talking about all genres . I'm an active OST listener since the beginning of the N64/PS1-era to this generation (and upcoming if you count Vita/3DS/Wii U games). I've always liked Nintendo music but most other companies like Sony/Microsoft and smaller game developers seem to go a similar way: going to a movie-like experience and therefore all their music sound the same. Of course, there are some exceptions like God of War and the first Halo-games. But most are losing their thematic tunes and are going into an orchestral cinematic experience. Nintendo has music you never hear anywhere else of almost any kind. Epic, fun, happy, calm, beautiful and simple but clever. So I've been listening almost solely the OST of Nintendo games the last few years. Especially Xenoblade and Super Mario Galaxy are extraordinary and instant classics (also the games themselves). I'm now enjoying the Mario&Luigi Dream Team OST a lot. My favorites are listed below.
Do you guys agree, or are there some beautiful soundtracks from Sony/Microsoft I should try?
Of course, tastes are pretty much subjective. But I find that all three companies have enjoyables OSTs.
Xenoblade and M&L OST are by a 3rd party composer.
Nintendo hires composers that can create memorable soundtracks (Koji Kondo (Mario, Zelda), Kenji Yamamoto (DKCR, Metroid), Mahito Yokota (Galaxy), and Hirokaza Tanaka (Animal Crossing) did the games you listed, and I also like Yuka Tsujiyoko (Fire Emblem), Jun Ishikawa (Kirby) and Junichi Masuda (Pokemon).
But there are others in the industry that can do this. Yoko Shimomura who worked on Dream Team and Xenoblade that you listed has plenty of other games. Also look at Motoi Sakuraba (Tales of), Nobuo Uematsu (Final Fantasy) and Yuzo Koshiro (Ys).
It's more like, only Japan can do it. As you say, the West are more interested in background, film score like composers. Again though, there are a few who are capable of being memorable. Jeremy Soule (Skyrim) is someone who does both.
Lately, I feel games are ashamed of BEING games and all want to be cinematic. From Versus XIII to Skyward Sword to Starcraft II, everything wants to be a film now. So memorable soundtracks are getting rarer. Plus all of the 80s composers are retiring/moving away from composing.
Many Nintendo sountracks give a good feel for the game im playing, and they are varied. Koji Kondo is a freaking genius.
I think Halo games have pretty good Ost but the most memorable are usually the main themes.
I thought Skyward Sword OST was indeed a bit dissapointing, but I can't say it was cinematic? Do you have some examples?
I only listed 2 Xenoblade songs but believe me I've more than 20 and they're all beautiful. Never expected a modern game to have such beautiful music. It feels new but also retro JRPG-like.


This entry was posted on 26 February 2017 at 7:30 pm	and is filed under Solving Blogs. You can follow any responses to this entry through the RSS 2.0 feed. You can leave a response, or trackback from your own site.
Great Setter's Blog – thanks Aragon!


"Weight" and other ceramic sculptures by Jessica Gardner will be on display at the Manassas Campus December 4 to January 31.
Jessica Gardner is displaying her ceramic sculptures in the Manassas Campus Colgan Hall Art Gallery through January 31. A reception to meet the artist will be held in the gallery space on January 22 at 12:00 p.m.
Gardner uses ceramic figures and interactive toys in her sculptures to trigger playful memories. She also incorporates second-hand children's clothing.
Gardner's work has been shown throughout the United States and published in "500 Teapots, Volume 2." She earned a Master of Fine Arts degree at Georgia State University and a post-baccalaureate certificate in ceramic arts at the University of Florida.
Gardner is the Annandale Campus coordinator of student engagement and also teaches studio art courses at the Manassas Campus.
For more information, contact Professor Barbara Lash at blash@nvcc.edu.


text:
Q: how to use axios on nativescript-vue hello i am trying fetch data from my backend but its look like not working also can't debug whats happening on load
i did try something like this
import axios from 'axios/dist/axios';

axios.get('http://api.chillwisper.tk/api/videos')
.then(response => {
  console.log(response.data);
  this.videos = response.data.data
})

nothing happen no console logs nothing
i also tried this
fetch("http://api.chillwisper.tk/api/videos")
            .then((response) => {
                console.log(response.data);
                console.log('this is test shit')
                this.videos = response.data.data
            })
            .catch((e) => {
                console.log(e.response)
            });

still not work how can do this i am new in nativescript but i know vue


Three public high school students present their original play about the Roman-era decision to celebrate the birth of Jesus on December 25, much to their teacher's chagrin.
Four characters; 3 men, 1 woman.
Approxomately 15 min. run time.
Commissioned in 2009 by EP Theatre.
This entry was posted on December 29, 2009, 7:55 pm and is filed under Plays, Short Plays. You can follow any responses to this entry through RSS 2.0. You can leave a response, or trackback from your own site.


text:
Julia Hartley-Brewer: 'We can fund public sector pay by cutting ridiculous foreign aid budget'
Julia Hartley-Brewer says the UK should cut foreign aid
Julia Hartley-Brewer has called for foreign aid to be cut in order to fund increases in public sector pay.
She made the comments during an interview with Conservative MP Nigel Evans, as they spoke about the growing calls on Prime Minister Theresa May to lift the public sector pay cap.
The talkRADIO presenter said: "I accept the whole 'charity begins at home' thing. I'm very much of the view that we have a duty to help out people in poorer nations.
"[But] it is quite clear by now, surely, that £13 billion is an absurd amount of money. No other Western countries are spending anything like this sum of money, this percentage of their GDP, it's a nonsense - it's got to stop.
"If you half that overnight you could pay for all of the pay rises in the public sector you want to."
Listen to the audio above
Public sector pay cap
Nigel Evans

redpajama_set_name: RedPajamaCommonCrawl


No one will argue that increasing the value of your home through home improvement projects is a great idea. However large home improvement projects can become quite expensive. Home improvements lighten your wallet and empty your savings account. Careful planning and thinking about all your financing options is necessary before beginning your home improvement project. Below are a few tips for home improvement home equity loan financing to take into consideration.
Home improvement home equity loans are great loans for home improvement, because the project can greatly increase the appraisal value of your home. This is a loan that is obtained to be able to get additional investments for use in the future. Home improvement projects such as bathroom additions, bedrooms and home extensions can increase the value of a house. However, some home improvement projects don't really result in increasing the value of the house. The construction of a swimming pool is one such project.
Take care when getting a home-improvement home equity loan. Don't forget that the collateral that you are putting up against the loan is your own house. If you can't make the payments and make them on time, you could end up losing your home. You borrowed money for the sole purpose of improving your house and losing your house would be a disastrous situation indeed.
Many people use home improvement home equity loans for other reasons. The money is sometimes spent finance other expenses such as vacations or everyday needs. The steady appreciation of their houses is what people rely on to be able to pay for the debt. If the value of their house depreciates at the end of any period, they are in huge financial hot water. This is why home improvement home equity loans should be used for the improvement of your home because the risks of depreciation are lower.


2018 a season of "BIG BUCKS"?
Discussion in 'Iowa Whitetail Conference' started by Sligh1, Aug 8, 2018.
So- talking to basically anyone I know- this season is far better than last. On average. Any place in the state, any neighborhood on average. Way better than 2017. Yes- u have neighborhoods that got shot to heck last year or didn't get rain this year, whatever. BUT.... on average the big deer folks have is far better based on 90% folks across state I know.
Are you guys seeing this BROADLY & ON AVERAGE???
1) drought last year destroyed huge growth potential on majority of farms. This year instead of massive drought- most of the state has had good rain. Reasons rain & lack of stressful massive heat are nutrition & massive browse amounts, plenty of water, no stressed & heat exhausted deer, etc. biggest reason IMO is the "better" amount of rain most got vs last year.
2) pretty much recovered from ehd. 2012 & 2013 were worst on average. I got PUMMELLED in 2015. BAD!!! Well- that's 5 years needed (or even more for quantities of deer to catch up) to really replenish Age structures. Most areas are out far enough from EHD this is actually the year where things genuinely are "caught back up" from the devastation. In reality, could be a couple more years of improvement still if we get lucky enough to avoid massive ehd outbreaks. Each year out from devastation has been improvement though.
3) Much like ehd.... being 2-3 years out FINALLY from the LATE SHED BUCK SEASON (RIFLES TO LATE JANUARY!) also is bearing the fruits of a replenished and "healed" deer herd that sustained much stress, damage & degradation to deer herd. Finally recovered or continuing.
My fear is: 1) EHD cycles again. 2) Dnr expands their shed buck season again with excuse of cwd. Step by step re-implementing it. Citizens have a "little" control on 1st one and total control on 2nd one. 2nd one is far messier with special interests lobbying behind scenes & hunters who don't or won't get involved & shrug it off. 3) further liberalization of regs, new weapons, new seasons, changing deer regs period that put more pressure & killing on the fragile resource.
Anyways, 95% of this post is positive & wondering if others are seeing same? We really look like we are heading into a "much better" 2018 season STATEWIDE. Folks will still deal with access issues, hard hunted areas, tough hunting BUT, IMO- over most the state- this looks MUCH BETTER than 2017. General thoughts or feedback?
**i have my eye on buck I call Chad 2.0!! . 5 years old and he's all of 110"!!!!! . Ha!!
Worried about EHD down south again this year. All the rain at the beginning then very little throughout the summer. Could be seeing outbreaks soon as more and more midge habitat is forming along the pond edges and creek beds as they dry. Big antlers having a bigger surface area for the midge to bite. Hope this isn't what happens but it's shaping up to have an outbreak again.
On another note have seen quite a few deer blow up in size this year from last year so hopefully they stay healthy.
4DABUCKS and Ishi like this.
The EHD stuff has been on my mind also the last month. I've talked to many deer hunters lately and everyone has the same sediments.
It wasn't drought last year alone IMO. SW IA and NW MO had good rain and near record crops yet antler growth was terrible. I hope you all are right for 2018 as some of my less than stellar performers in antler growth from last year made it to 2018 and should be another year older. We have had heat and drought stress around here this year though. Haven't saw them recently but found their sheds so we will see.
It seems like I have seen more bigger bucks this year so far.
2017 was the year of the "shrinking deer". All over, multiple farms. 2018 is looking good, although it did appear to me they got off to a slow start and are finishing growing later. I've seen more deer in general than ever before. This is not necessarily a good thing but it has been noticeable. Does for miles.
I must have an odd ball farm, last year almost every buck on my place 3 years or older made at least a 10-15 inch jump some of the young guys made 40-50 inch jumps. Now this year every buck I have over 5 is either the same size or got smaller, they also seem to be way behind on growth compared to last year.One buck in particular was pretty much finished this time last year and this season he is completely balled up and barely has g4s, anyone else seeing anything similar?
deep woods goat hunter, YtailFreak, NWBuck and 4 others like this.
My cams are showing fewer bucks, smaller on average than previous years. Usually I get at least one that grabs my attention, sometimes two. Nothing so far this year. I need to get the rest of my cameras out, been too busy.
In the past 5 years this is shaping up to be one of the worst years by far. I dont have anything on camera so far that's worth getting excited about. We've gotten some rain recently but it's been real dry this year, unfortunately seems like prefect conditions for EHD again. Not just bucks but our doe population is low, we may not be killing any again this year. I'm praying something worthwhile moves in after the velvet comes off.... Putnam Co MO.
I totally get that. Here is where it's interesting for me too.... for one farm I hunt, I shot a 215-ish on it and considered a "good farm". I don't have any giants there. Nothing "oh my gosh" at all. It's funny cause my buddies would think I automatically would. NOPE. My hope on even hunting some "absolute giants" which is rare anywhere - will be on smaller farms that no one would think "that screams giants". Every particular farm can have tough years - even while others, on average, are outstanding. I know exactly what you are dealing with or I've had your sentiment before.
On average, driving around to seeing pictures folks I know have- seems way up. Why I actually will spend a lot of time on smaller farms because their average on bucks & size is up a good bit.
"THE AREAS FULL OF BIG BUCKS.... DON'T REALLY HAVE MANY BIG BUCKS". So many folks think "if i owned this, if I hunted HERE, if I was THERE- be tons of giants!!!" & on certain years- any random farm or big farm CAN have some serious slammers. On average, it's usually still very difficult and many years, any farm u could dream up- will have extremely challenging hunting. Even years where you would say "nothing giant period". Just a little side rant. .
I'm trying to stay optimistic, I know anything could walk thru at any time. And I do think I'm in a "pocket" where we've taken a down turn and it should bounce back eventually. Just several miles away people I talk with are very happy. With all the habitat work I'm doing these dang bucks don't know what they're missing!!
I can't form an opinion, yet. The flood waters finally receded, and the deer are starting to resume their normal haunts. Hopefully some deer I have some history show back up in the next few weeks so I have some idea.
Too early to tell for me. Haven't had cameras out very long, but talk from my dad and grandpa about numbers and size that they've seen is typical of any other year.
I've got nothing... but there's so much ag/draws around me I dont normally pick up the deer I hunt in the fall until velvet switcheroo. When the beans turn/corn is cut the WTD's that I'm interested in tend to show up around 10/1.
My urban spot is kind of the same, but not much crops around. They just don't show up really until October. Which I am fine with.
Im in NCMO, Randolph county.


text:
> Where Are They Now
> World Championship team bosses
> John Clarke
John Clarke
No picture available of John Clarke as a Team Boss.
No current picture available of John Clarke.
Lambeth, London
Williams (1973-1974)
A quiet, reserved man, who was formerly at March, he worked on the Iso-Williams project and in Formula 1 until the late 1970's, as well as designing the Boxer Formula 2 car, but has been little heard of lately. His spell with Williams came to an unhappy end due to a disagreement with Frank Williams as, in John's own words - there had never been any money for development (nor quite often for wages) and I therefore felt a bit of a spare part and was keen to return to designing things which would actually get made - not, as is sometimes reported, due to the involvement of Dallara in the car's development. Clarke, away from racing, ran his own firm, JCM Engineering initially in Surrey and now in Hampshire. John remains involved, even in his seventies. He also ran an engineering support company in Alton, Hampshire.
Biography last updated 16 Jan 2020
team boss
Other World Championship team bosses from Great Britain
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Other World Championship team bosses born in 1944
Akira Akagi, Rory Byrne, Matsuhisa Kojima, Harvey Postlewaite.