although cdte - based czt 
 ( cd1-xznxte ) crystals are the leading semiconductor 
 compounds for radiation - detection applications , the widespread deployment of czt 
 detectors has been limited by the high cost ( due to a low manufacturing 
 yield ) of detector - grade materials .
property nonuniformity has been 
 the major cause for both poor performance and a low yield of usable 
 portions of ingots .
it has long been established 
 that micrometer - scale defects such as tellurium inclusions / precipitates 
 affect carrier transport and uniformity .
however , extensive previous efforts to reduce te 
 particles have not resulted in an anticipated improvement in czt property 
 uniformity .
one special structural feature of czt crystals is their 
 extreme softness and their propensity to always develop a high density 
 of dislocations during synthesis .
experiments have shown 
 that when these dislocations are organized into a network of subgrains , 
 they can directly affect charge carriers to cause electric 
 field perturbations .
clearly , 
 eliminating or at least controlling dislocation subgrain structures 
 holds great potential to further improve czt crystals .
many 
 types of dislocations may be present in czt crystals , including 
 shuffle dislocations , glide dislocations ,  and  dislocations , 
 partial and perfect dislocations , edge , screw , mixed dislocations , 
 and misfit dislocations . the understanding of the behavior of the 
 various combinations of these dislocations in conjunction with developing 
 methods for improving the corresponding dislocation network structures 
 are therefore challenging . the use of predictive , high - fidelity atomistic 
 simulations can help screen out these structure - limiting dislocation 
 types . here
we use a recently developed czt bond order potential ( bop ) in lammps ( large - scale atomic / molecular massively parallel simulator ) to perform selected molecular dynamic ( md ) simulations to determine 
 configurations , energies , and mobilities of 29 different types of 
 dislocations in cdte crystals .
[ the czt ternary bop is an extension of the cdte binary bop , so they are equivalent if used only for cdte .
] an efficient 
 method to derive activation free energy and activation volume of thermally 
 activated dislocation motion will be explored .
our focus is to identify 
 and validate important dislocations and dislocation phenomena in the 
 material , which can reduce future experimental efforts to improve 
 the dislocation network structures .
the key to high - fidelity 
 md simulations of dislocations is an interatomic 
 potential that is transferrable to the dislocated regions where the 
 local atomic environment is significantly different from that of the 
 perfect bulk crystal .
details of the czt bond order potential applied 
 here are omitted as they are very complex and have been fully described 
 previously .
it is demonstrated that the czt bop captures well 
 the property trends of a variety of elemental and compound structures / phases 
 ( with local environment varying from coordination 2 to 12 ) , simulates correctly the crystalline growth of stoichiometric compounds 
 during both vapor deposition ( under nonstoichiometric vapor fluxes ) and melt - growth processes , and predicts precisely the configuration and density 
 of misfit dislocations in vapor deposited multilayers .
note that the correct simulation of growth , especially 
 under a nonstoichiometric environment , is a very strong validation 
 of the transferability of a potential as the process involves chemical 
 reactions and samples a variety of configurations and chemistries 
 statistically formed on a growth surface .
for this reason , the czt 
 bop is expected to produce reasonable results regarding dislocations .
different 
 types of dislocations can be constructed on various planes 
 in a zinc - blende crystal .
for instance , mobile dislocations usually 
 lie on the { 111 } slip planes , and misfit dislocations in multilayered 
 films are often seen to lie on interfacial planes , which are not necessarily parallel to { 111}. here we only 
 study dislocations that are on the { 111 } slip planes , with burgers 
 vector b being either partial < 112>a/6 
 or perfect < 110 > a/2 .
( a ) 
 front projection of 
 ( 112 ) planes , ( b ) top projection of two ( 111 ) planes a and 
 b showing glide perfect dislocations , ( c ) top projection of two ( 111 ) 
 planes a and b showing glide partial dislocations , and ( d ) top projection 
 of two ( 111 ) planes b and c showing shuffle dislocations .
( a ) front projection 
 of 
 ( 110 ) planes , ( b ) top projection of two ( 111 ) planes a and 
 b showing glide perfect dislocations , ( c ) top projection of two ( 111 ) 
 planes a and b showing glide partial dislocations , and ( d ) top projection 
 of two ( 111 ) planes b and c showing shuffle dislocations . assuming that the dislocation line direction  aligns with the z direction 
 and the slip occurs 
 on the x  z plane , our definition 
 of different types of dislocations is shown in figures 1 and 2 , where dark and light circles 
 distinguish metal ( cd or zn ) and te atoms , respectively .
dislocations 
 defined in figure 1 are based upon an edge 
 dislocation coordinate system shown in figure 1a where the dislocation line  is 
 perpendicular to one of the burgers vectors < 110>a/2 ( x axis ) .
dislocations defined in figure 2 are 
 based upon a screw dislocation coordinate system shown in figure 2a where  is parallel 
 to one of the perfect burgers vectors < 110>a/2 ( z axis ) .
figures 1a and 2a indicate that a dislocation slip can occur at two different locations , 
 one between two closely separated planes a and b ( where the interplane 
 bonds are not perpendicular to the planes ) , and the other between 
 the more widely separated planes b and c ( where the interplane bonds 
 are perpendicular to the planes ) . according to convention
, dislocations 
 between a and b are called glide dislocations , and dislocations between 
 planes b and c are called shuffle dislocations .
burgers vectors 
 of glide dislocations can be examined from a top 
 projection of two isolated planes a and b. figures 1b and 2b use such a projection to explore 
 perfect dislocation burgers vectors ( < 110>a/2 ) .
it can be seen 
 from figure 1b that in the edge coordinate 
 system , there are three nonequivalent perfect burgers vectors that 
 can be represented by the angle  between b and  , i.e. ,  = 30 , 
 90 , and 150. there are other perfect burgers vectors , 
 but they are equivalent to the ones listed due to crystal symmetry . 
 for instance , the  = 330 dislocation is equivalent to 
 the  = 30 dislocation .
similarly , it can be seen from 
 figure 2b that in the screw coordinate system 
 there are three more nonequivalent perfect burgers vectors  
 = 0 , 60 , and 300. figures 1c and 2c 
 use the same top projection of two isolated planes a and b to explore 
 partial dislocation burgers vectors ( < 112>a/6 ) .
it can be seen 
 from figure 1c that in the edge coordinate 
 system , there are two nonequivalent partial burgers vectors  
 = 0 and 120. similarly , it can be seen from figure 2c that in the screw coordinate system there are 
 two more nonequivalent partial burgers vectors  = 90 
 and 330. burgers vectors of shuffle dislocations can be 
 examined from the 
 top projection of two isolated planes b and c. figures 1d and 2d use such a projection to explore 
 shuffle dislocation burgers vectors . here the light and dark atoms 
 in the b and c planes are aligned in the projection ( i.e. , only light 
 atoms are seen ) , and hence , there are no partial dislocation burgers 
 vectors for shuffle dislocations .
as indicated in figure 1d for the edge coordinate system , we explore three shuffle 
 dislocation burgers vectors  = 30 , 90 , and 150 , 
 which are the same as those in the perfect glide dislocation case 
 shown in figure 1b .
as indicated in figure 2d for the screw coordinate system , we explore three 
 more shuffle dislocation burgers vectors  = 0 , 60 , 
 and 300 , which are the same as those in the perfect glide dislocation 
 case shown in figure 2b .
dislocations 
 shown in figures 1 and 2 always occur between a metal ( cd or zn ) and te 
 { 111 } planes .
if a dislocation has an edge component , then the dislocation 
 core can have extra metal atoms or extra te atoms . according to convention , 
 a dislocation with more metal atoms at its core
is called an  
 dislocation , and a dislocation with more te atoms at its core is called 
 a  dislocation . according to the discussion above
, we 
 will consider a combination 
 of glide and shuffle dislocations , partial and perfect dislocations , 
 various burgers vectors , and  and  dislocations .
a cdte crystal 
 system with 336 ( 112 ) planes in the x direction , 
 60 ( 111 ) planes in the y direction , 
 and 28 ( 110 ) planes in the z direction is 
 used to calculate the ( 111 ) stacking fault energy sf .
the stacking fault is created by shifting the upper part of a perfect 
 crystal by a partial dislocation vector with respect to the lower 
 part ( at a glide dislocation slip plane between planes a and b , see 
 figure 1 or 2 ) .
molecular 
 statics ( ms ) energy minimization simulations are used to calculate 
 relaxed energies of both the perfect and the stacking faulted crystals 
 under the periodic boundary conditions in both x and z directions 
 and a free boundary condition in y direction . the stacking fault energy 

is then calculated as sf = ( esf  e0)/(lxlz ) , where esf and e0 are respectively the relaxed energies for stacking faulted 
 and perfect crystals and lx and lz are respectively 
 the system dimensions in x and z directions .
note that the surface energy in the y direction does 
 not affect the calculation as it exists for both stacking faulted 
 and perfect crystals and therefore cancels out .
cdte 
 is known to have a low stacking fault energy around 10 mj / m .
this in turn allows dislocation energies , i.e. , the 
 energy difference between a dislocated crystal and a perfect cdte 
 crystal containing the same number of atoms , to be calculated as a 
 function of dislocation core radius .
the key for the calculations 
 is to use the initial dislocation configurations that can be fully 
 relaxed during md and ms energy minimization simulations .
we find that 
 an initial dislocation created by simply positioning atoms according 
 to the displacement field of the linear elastic continuum solution 
 does not always fully relax at the core . instead
, a dislocation created 
 by shearing the relevant portions of the crystals using md simulations 
 is found to normally relax to lower energies than the other methods . 

in particular , as will be described below , an md scheme illustrated 
 in figure 3 is effective in creating partial , 
 undissociated perfect , and shuffle dislocations , and an md scheme 
 illustrated in figure 4 is effective in creating 
 dissociated perfect dislocations .
these schemes enable atom positions 
 to be strictly determined from atom interactions and also best reflects 
 realistic scenario that dislocations are created under local shear 
 stresses . line energy model for partial , undissociated perfect , and shuffle 
 dislocations .
( a ) md simulations to create dislocations and ( b ) md 
 and ms simulations to relax dislocations .
the computational system used is periodic in the x and z axes with free surfaces along the y axis . for the undissociated dislocations ( including partial 
 and shuffle dislocations ) shown in figure 3 ,
the system is divided into different regions as indicated by blue 
 ( middle upper ) , orange ( middle lower ) , black ( left and right ) , and 
 white ( the rest ) colors . with the black regions held fixed and the 
 white regions free to relax , an md simulation is first performed to 
 uniformly move the blue region by half of the burgers vector 0.5b and the orange region by negative half of the 
 burgers vector 0.5b. for partial 
 dislocations , the movement of the blue and orange regions is completed 
 directly over a total simulated time of 0.1 ns . for perfect dislocations , 

the movement is through two stages over a total simulated time of 
 0.2 ns to follow the lowest energy path of two partials , i.e. , 0.5(b1 + b2 ) ( = 0.5b ) for the blue region and 0.5(b1 + b2 ) ( = 0.5b ) for the orange region . 

in addition , the system temperature is uniformly decreased from 300 
 to 10 k during simulation .
this md procedure creates two dislocations 
 with opposite burgers vectors ( also note that if the dislocations 
 have an edge component , then one is an  and the other one is 
 a  dislocation ) as shown in figure 3a . for the dissociated perfect dislocations shown in figure 4 , the system is divided into different regions as 
 indicated by light and dark blue ( middle upper ) , light and dark orange 
 ( middle lower ) , black ( left and right ) , and white ( the rest ) colors . 
 with the black regions held fixed and the white regions free to relax
, 
 an md simulation is first performed to uniformly move the light and 
 dark blue region by half of a shockley partial burgers vector 0.5b1 and the light and dark orange region 
 by negative half of the same shockley partial burgers vector 0.5b1 over a total simulated time of 0.1 
 ns . a second md simulation
is then performed to move the dark blue 
 region by half of the other shockley partial burgers vector 0.5b2 and the dark orange region by negative 
 half of the same shockley partial burgers vector 0.5b2 over another total simulated time 
 of 0.1 ns .
in addition , the system temperature is uniformly decreased 
 from 300 to 10 k during these two simulations .
this md procedure creates 
 the same two perfect dislocation burgers vectors as shown in figure 3a except that each perfect dislocation is dissociated 
 into two shockley partials bounding a stacking fault band as shown 
 in figure 4a .
( a ) md 
 simulations to create dislocations and ( b ) md and ms simulations to 
 relax dislocations .
once dislocations are created in the md simulations , 
 the blue and 
 orange regions are released . a narrow region in the middle of the 
 system
is allocated as shown by the black color in figures 3b and 4b . for perfect dislocations , 

this narrow region remains a perfect crystal because the blue region 
 is moved by a perfect burgers vector with respect to the orange region . 
 for partial dislocations ,
the narrow middle region contains a stacking 
 fault . with left , right , and
middle black regions fixed , an ms simulation 
 is then conducted to relax the dislocations . in this work ,
we 
 use an edge dislocation geometry containing 97 
 ( 110 ) planes in the x direction , 60 ( 111 ) 
 planes in the y direction , and 24 ( 112 ) planes 
 in the z direction and a screw dislocation geometry 
 containing 168 ( 112 ) planes in the x direction , 
 60 ( 111 ) planes in the y direction , and 14 ( 110 ) 
 planes in the z direction to create various perfect 
 and dislocated computational crystals .
relaxed energies of both dislocated 
 and perfect crystals are calculated . to separate  and  
 dislocation effects , the total energy is split to the left and right 
 parts respectively by summing up the atomic energies of the left and 
 right halves of the systems .
the line energies of the left and the 
 right dislocation , i , are then 
 calculated as i = ( ed , i 
wilzsf)/lz , where ed and e0 are the 
 relaxed energies of the dislocated and the perfect crystals , respectively , lz is the dislocation length 
 (= system dimension in the z axis ) , sf is the stacking fault energy , w is the stacking 
 fault bandwidth , and the subscript i (= l or r ) indicates 
 the left and the right dislocations .
we find that of all the simulations using 
 the initial perfect dislocations , the maximum splitting distance is 
 only 25  after energy minimization .
table 1 indicates that the screw (  = 0 ) dislocations 
 have lower energies than other dislocations , being 0.38 , 0.69 , 0.79 , 
 and 0.81 ev /  for partial , dissociated perfect , undissociated 
 perfect , and shuffle dislocations , respectively .
interestingly , the 
 acute angle of the second lowest energy ( shuffle , glide partial , glide 
 perfect ) dislocations all equal 30 ( e.g. ,  = 30 , 
 150 , and 330 ) . in addition , the energies of dissociated 
 perfect dislocations are about 8092% of the energies of undissociated 
 dislocations , suggesting that the dissociated dislocations are more 
 stable .
however , the observation that the perfect dislocations remain 
 approximately undissociated during md and energy minimization simulations 
 suggests an energy barrier of dissociation .
the configuration of perfect 
 dislocations will be examined further below in md simulations of dislocation 
 motion .
( a ) glide 
 partial dislocations , ( b ) glide perfect dislocations , and ( c ) shuffle 
 dislocations .
the dislocation line energy is not a constant material 
 property , 
 but rather increases with the material volume within a radius r from the dislocation core .
the dislocation line energy 
 at a given radius r can be calculated using the same 
 approach described above except that the energies of the perfect and 
 dislocated crystals are taken as the sum of the atomic energies for 
 all the atoms within the radius .
the dislocation line energies as 
 a function of r thus obtained are shown in figure 5a 
figure 5a  c indicates that dislocation 
 line energies calculated with the atomistic models linearly increase 
 with dislocation radius r when r is larger than 40  .
this is consistent with the boundary conditions 
 used in figures 3 and 4 where the black regions are held fixed during simulations .
the dislocation 
 core energies at small radii , say r < 30  , 
 come from the relaxed dislocation core structures . in the present 
 work , the system sizes have been chosen to be sufficiently large so 
 that the predicted dislocation core structures are independent of 
 the boundary conditions . as a result
, our simulations can give accurate 
 dislocation core energies at r < 30  that 
 can not be predicted from linear elastic theories .
it can be seen from 
 figure 5a  c that dislocation energies 
 sharply decrease when the radius is decreased .
in addition ,  
 and  dislocations have similar energy vs r curves for partial 
 and shuffle dislocations . while  and  dislocation energies 
 are also similar for the perfect dislocations at large radii , they 
 become different at the core .
in particular , the  perfect dislocation 
 core energies are lower than the  perfect dislocation core 
 energies .
one effective approach to simultaneously study the motion of multiple 
 dislocations is to examine the evolution of a dislocation loop under 
 a shear stress .
the computational system consists of a cdte crystal 
 with 120 ( 112 ) planes in the x direction , 
 18 ( 111 ) planes in the y direction , and 68 ( 110 ) 
 planes in the z direction .
periodic boundary conditions 
 are used in the x and z directions 
 and a free boundary condition is used in y direction 
 during the simulation . to create a dislocation loop ,
the system is 
 divided into different regions as shown in figure 6(a ) : the black ( top and bottom ) regions , the blue ( middle 
 upper ) hexagonal region , the orange ( middle lower ) hexagonal region , 
 and the white ( remaining ) region . as plastic deformation usually proceeds 
 through a glide of dislocations
hence , 
 we assume that the boundary between the upper and lower hexagons is 
 between planes a and b as defined in figures 1a and 2a . to create a realistic dislocation 
 loop with a burgers vector of b ,
an md simulation 
 is performed to uniformly move the top black and blue regions by 0.5b and the lower black and orange regions by 0.5b over a simulated time of 0.1 ns while allowing 
 the white region to relax according to newton s equation of 
 motion and a constant temperature of 10 k. as mentioned above , for 
 a partial dislocation loop , the md simulation is completed directly , 
 and for a perfect dislocation loop , the md simulation is completed 
 in two stages to follow the low energy path of two partial dislocations . 

this md simulation procedure creates a dislocation loop at a shear 
 strain of approximately  = b / t , where b is the magnitude of the dislocation burgers 
 vector and t is the thickness of the sample . the 
 system is then redivided into three regions as shown in figure 6b : the black ( top and bottom ) regions and the white 
 ( remaining ) region .
the evolution of dislocation loops is examined 
 in subsequent md simulations at a high temperature of 900 k where 
 the top and bottom surface black layers are further moved in positive 
 and negative burgers vector directions to increase the shear strain .
( a ) md simulation to create dislocation 
 loop and ( b ) md simulation to evolve dislocation loop . a partial 
 dislocation burgers vector of [ 112]a/6 is used to create an 
 initial dislocation loop using the procedure shown in figure 6(a ) .
further shear strain is applied 
 by moving the top and bottom black regions in the + x and x directions by 1.394  , respectively , 
 using the procedure shown in figure 6b .
the equilibrium dislocation loop configurations obtained 
 at 900 k temperature are examined in figure 7 by showing the top view of four consecutive planes above and below 
 the slip plane , where blue and gray circles indicate cd and te atoms 
 respectively , and the yellow hexagon indicates the initial partial 
 dislocation loop at xy = 0.039 . 

panels a and b in figure 7 are obtained using 
 exactly the same conditions except that the cd and te atoms are switched 
 between the two figures , so that in figure 7a , the right side of the loop is an  dislocation and the left 
 side of the loop is a  dislocation , whereas in figure 7b , the left side is  and the right side is 
 . because the movement of a partial dislocation leaves 
 behind a stacking fault , the relaxed partial dislocation loop can 
 be identified by the boundary of the two stacking patterns demonstrated 
 in the figure .
it can be seen from figure 7a that the hexagonal dislocation loop becomes asymmetric under the 
 shear and in particular , the  dislocation moves by a far longer 
 distance than the  dislocation .
this is further verified in 
 figure 7b where the asymmetry of the dislocation 
 is flipped as a consequence of the flip of  and  dislocations . 

this simulation strongly indicates that  dislocations are much 
 more mobile than  dislocations ( despite their similar line 
 energies in table 1 ) , in good agreement with 
 experiments conducted for many semiconductor compounds .
( a ) configurations 
 at strains xy = 0.039 and xy = 0.078 and ( b ) the same as in panel a 
 except that cd and te atoms are switched
t = 0.0025 ( a ) , 0.0100 ( b ) , 0.0250 ( c ) , and 0.0350 ns ( d ) .
one interesting observation in figure 7 is 
 that to accommodate the motion of the  segment (  = [ 110 ] ) of the dislocation loop , two segments (  = [ 112 ] ) are created .
this means that the motion 
 of the mobile  partial dislocations can cause elongated screw 
 partials . as a result
, cdte crystals may have a high density of such 
 type of screw partials , in agreement with its low energy shown in 
 table 1 . on the other hand ,
this corresponds 
 well to the high energy of the edge partial in table 1 .
a perfect 
 dislocation loop with a burgers vector of [ 110]a/2 is created 
 using the procedure described in figure 6a . 

this procedure produces an initial shear strain of yz,0 = 0.068 . by moving the top and bottom black regions 
 in the + z and z directions 
 at a constant speed of 1.2 nm / ns during an md simulation at 900 k , 
 as described in figure 6b , the system is further 
 deformed at a strain rate of dyz / dt = 0.338/ns .
the configuration of the dislocation 
 loop during the further strain is shown in figures 8a  d as a function of time , where the color scheme represents 
 the magnitude of the slip vector developed previously to show dislocations .
figure 8a shows that 
 shortly after further strain at t = 0.0025 ns ( yz = 0.069 ) , the dislocation loop remains 
 hexagonal but some portion of the  dislocation ( the lower part 
 of the [ 101 ] section ) bows out .
figure 8b shows that at t = 0.0100 ns ( yz = 0.071 ) , the dislocation bows out more . in figure 8c at 0.0250 ns ( yz = 0.076 )
, the  dislocation is seen to have moved a significant 
 distance that causes it to pass through the low boundary and enter 
 into the image from the top boundary under the periodic boundary condition . 

clearly , perfect  dislocations are also much more mobile than 
 perfect  dislocations ( despite their similar line energies 
 shown in table 1 ) , in good agreement with experiments . interestingly
, figure 8c indicates that , although screw dislocations also have higher mobility 
 than  edge dislocations , they are much less mobile than  
 edge dislocations . when the  dislocation meets the  
 dislocation at the top of the hexagon , they are annihilated , and the 
 dislocation loop becomes two vertical screw dislocations running across 
 the periodic dimension in the z direction , as can 
 be seen in figure 8d at t = 
 0.035 ns ( yz = 0.080 ) .
further 
 strain will cause the two dislocations to continuously move in the 
 x directions until they completely sweep the 
 entire area and are annihilated .
interestingly , we find that the dislocation 
 motion is thermally activated and does not occur when the simulation 
 is performed at a low temperature , say , 10 k. in addition , we find 
 that the dislocations do not move as straight lines , but rather through 
 the formation and rapid motion of kinks along the z direction as shown in figure 8c .
this is 
 consistent with the traditional kink model of thermally activated 
 motion of dislocations .
the observation 
 that screw dislocations are much less mobile than 
  edge dislocations indicates that when the  sector of 
 a dislocation loop moves , the screw dislocation sectors (  = [ 110 ] ) will elongate . as a result
this is consistent with the low 
 energies of screw dislocations shown in table 1 .
in addition , no dislocation dissociation is observed in figure 8 , indicating that the perfect dislocations do not 
 necessarily dissociate under nonequilibrium conditions .
the details 
 of dislocation mobility are important to understand 
 because they directly control dislocation network structures . for 

thermally activated motion of dislocations , the dislocation velocity 
 can be expressed as1where q and  are respectively 
 the activation energy and activation volume of the dislocation motion , 
  is the shear stress , k is boltzmann s 
 constant , t is the temperature , and v0 is a constant representing saturation dislocation velocity 
 for barrierless ( very high temperature or shear stress ) dislocation 
 motion . here q and  are the primary parameters 
 quantifying dislocation mobility .
the discussion presented above 
 already indicates several distinctive 
 dislocation types that may determine dislocation structures .
first , 
  and  partial and perfect dislocations have a strong 
 edge component and can contribute to the climb motion responsible 
 for the formation of small angle grain boundaries .
 dislocations 
 are the most mobile , resulting in special structural features of dislocation 
 networks . screw partial and perfect dislocations can also be important . 

these screw dislocations are expected to have a high density as they 
 have low energies and can be created by the motion of  dislocations . 

the objective of the present work is not to exhaustively tabulate 
 the mobilities of a variety of dislocations .
hence , we focus on calculating 
 the activation energy barrier q and activation volume  for 
  and  edge type of perfect dislocations . such studies 
 will allow us to fully explore the method for calculating q and  and to validate the results using known experimental 
 relative mobilities of  and  dislocations .
in addition , 
 edge dislocations play an important role in dislocation network formation 
 due to their climb motion .
the extension of the method to other dislocations 
 that can provide complete inputs for simulations of dislocation network 
 structure evolution using larger scale ( e.g. , dislocation dynamics ) models will be performed in future work .
the activation energy barriers of dislocation motion are often 
 directly calculated from energy profiles using the nudged elastic 
 band method .
however , the stop - action observation of the dislocation migration 
 case shown in figure 8 indicates that the thermally 
 activated motion of dislocations proceeds through a variety of different 
 paths . without requiring an explicit treatment of migration paths , 

we perform md simulations of dislocation motion at a variety of temperatures 
 and shear stresses to fit an apparent activation energy barrier and 
 activation volume directly from ( 1 ) .
the activation energy determined 
 this way pertains to a free energy barrier that incorporates the multiple 
 path ( entropy ) effects .
in addition , this approach provides a direct 
 verification of ( 1 ) and , hence , the thermally activated dislocation 
 mechanism .
the initial crystal contains 144 
 ( 110 ) planes in the x direction , 36 ( 111 ) 
 planes in the y direction , and 24 ( 112 ) planes 
 in the z direction .
the system employs periodic boundary 
 conditions in x and z directions 
 and a free boundary condition in y direction .
first , 
 the initial crystal is divided into different regions as shown in 
 figure 9a : the black ( left and right ) regions , 
 the blue ( middle upper ) region , the orange ( middle lower ) region , 
 and the white ( remaining ) region .
since we only explore glide dislocations , 
 we assume that the boundary between the blue and orange regions is 
 between a and b planes as defined in figures 1a and 2a . to create dislocations with a perfect 
 burgers vector of b = b1 + b2 , an md simulation 

is performed where the black regions are fixed , the white regions 
 are relaxed , and the blue and orange regions are first uniformly moved 
 by 0.5b1 and 0.5b1 , respectively , over a simulated time 
 of 0.001 ns and then moved by 0.5b2 and 0.5b2 , respectively , 
 over another simulated time of 0.001 ns .
the crystal is then cut in 
 half in the x direction as shown in figure 9b with a corresponding adjustment of the periodic 
 length so that the two halves remain periodic in the x direction .
in particular , we use b = [ 110]a/2 , b1 = [ 121]a/6 , and b2 = [ 211]a/6 to create two crystals , one containing 
 an  perfect dislocation , and the other one containing a  
 perfect dislocation . edge dislocation mobility model . ( a )
md simulation to 
 create  
 and  dislocations , ( b ) isolate out  and  dislocations , 
 and ( c ) md simulation to evolve dislocation under shear stress xy . for dislocation motion simulations , the ( half ) 

system is divided 
 into three regions : the black ( top and bottom ) regions and the white 
 ( remaining ) region . by adding corresponding forces to the atoms in 
 the top and bottom regions in + x and x directions , respectively ,
molecular dynamics simulations 
 under an npt ( constant atom number , pressure , and temperature ) condition 
 are then performed to study migration of dislocations at a variety 
 of stresses and temperatures .
we find that in all of these simulations 
 the dislocation splitting is very small , further verifying the discussions 
 presented above that perfect dislocations do not necessarily dissociate . 

assuming that the slip vector reaches 
 a maximum at the dislocation core , the dislocation position is calculated 
 as a function of time .
results obtained for  and  dislocations 
 at selected stresses and temperatures are shown in figure 10 .
it can be seen from figure 10 that all the dislocation vs time curves approximately fall 
 on straight lines .
this shows that the data can provide reliable information 
 about the steady - state dislocation velocity .
the slope of these straight 
 lines changes with stress and temperature , indicating that the simulated 
 conditions capture the thermally activated mechanisms of dislocation 
 motion .
figure 10a shows the effect of 
 temperature 
 on dislocation motion at a constant shear stress of 0.5 gpa .
increasing 
 the temperature is seen to increase the dynamics , consistent with 
 the thermal activation mechanism . at the same temperature ,
the  
 dislocation is seen to move faster than the  dislocation , further 
 verifying the results seen in figures 7 and 8 .
figure 10b shows the effect 
 of shear stress at a constant temperature of 600 k. here increasing 
 the stress also correctly increases the dislocation velocity , and 
 the  dislocation is again validated to move faster than the 
  dislocation .
( a ) effect 
 of temperature 
 at a given stress and ( b ) effect of stress at a given temperature .
the steady state velocity can be easily calculated 
 from the slope 
 of the curves shown in figure 10 .
the results 
 obtained at a constant shear stress of xy = 0.5 gpa are plotted as ln(v ) vs ( kt ) curves in figure 11 . here
the unfilled and filled circles represent md data for 
  and  dislocations respectively , and the lines are fitted 
 to the data using equation ( 1 ) .
figure 11 indicates 
 that the md data falls fairly well on straight ln(v ) vs ( kt ) lines , validating 
 that the method indeed captures well the thermally activated dislocation 
 migration . through the fitting , we find that the activation energy 
 and activation volume are q = 0.14 ev and  
 = 17  for the  dislocation and q = 0.27 ev and  = 38  for the  dislocation . 

this again verifies that in the thermally activated regime ( e.g. , 
 xy  1 gpa ) , the  
 dislocation is more mobile than the  dislocation .
the 
 thermally activated plastic deformation of cdte has been experimentally 
 analyzed below room temperature . the 
 activation energy and activation volume
are assumed to decrease with 
 increasing stress in the experiments . under this assumption ,
the fitted 
 experimental activation energy and activation volume are approximately 
 0.8 ev and 5500  , respectively , at a small stress 
 of 0.005 gpa , and fall to < 0.5 ev and < 2250  , respectively , at a stress of 0.010 gpa . here
we assume that q and  are constant and are fitted using md data 
 at high stresses ( e.g. , 0.5 gpa ) .
in addition , simulations target 
 only the peierls potential barrier mechanism of specific (  
 and  perfect edge ) dislocations , whereas , experiments may include 
 point defect obstacle mechanisms and the contributions from other 
 dislocations . as a result
, it is expected 
 that the simulated activation energy and activation volume are smaller 
 than their experimental counterparts .
our bop - based md simulations 
 of dislocations in cdte crystals lead 
 to the following conclusions:(1) dislocations move much faster 
 than  dislocations , in good agreement with well - known experiments 
 for semiconductor compounds .
this , along with previous demonstrations 
 that our bop captures properties of many phases and predicts crystalline 
 growth during md simulations of chemical vapor deposition and melt - growth , provides strong validation 
 that the bop correctly captures the physics of dislocation motion 
 in cdte .
the bop - based md simulations , therefore , can provide high - fidelity 
 data to construct large scale ( e.g. , dislocation dynamics ) models that can simulate the evolution of dislocation 
 cell structures.(2)for 
 partial , perfect , and shuffle 
 sets , screw dislocations always have the lowest energies . dissociated 
 glide dislocations always have lower energies than perfect glide dislocations . 

however , perfect glide dislocations do not necessarily dissociate 
 due to the energy barrier of dissociation .
in particular , the splitting 
 distance of perfect glide dislocations is seen to remain very small 
 in our dynamic md simulations.(3)the acute angle of the second lowest 
 energy ( shuffle , glide partial , glide perfect ) dislocations all equal 
 30 ( e.g. ,  = 30 , 150 , 330).(4)screw dislocations are 
 much less mobile 
 than  edge dislocations . as a result ,
the screw dislocation 
 sectors perpendicular to the  sector are elongated when an 
  sector of a dislocation loop moves .
this suggests that screw 
 dislocations are likely to have a higher density than edge dislocations , 
 in good agreement with the low screw dislocation energies.(5)without calculating the 
 energy profile 
 explicitly , we show that md simulations of dislocation motion at a 
 variety of stresses and temperatures can be used to effectively deduce 
 the activation free energy barrier and activation volume of thermally 
 activated glide of dislocations .
in particular , we found activation 
 free energy barriers of 0.14 and 0.27 ev and activation volumes of 
 17 and 36  , respectively , for  and  
 edge types of glide dislocations .
 dislocations move much faster 
 than  dislocations , in good agreement with well - known experiments 
 for semiconductor compounds .
this , along with previous demonstrations 
 that our bop captures properties of many phases and predicts crystalline 
 growth during md simulations of chemical vapor deposition and melt - growth , provides strong validation 
 that the bop correctly captures the physics of dislocation motion 
 in cdte .
the bop - based md simulations , therefore , can provide high - fidelity 
 data to construct large scale ( e.g. , dislocation dynamics ) models that can simulate the evolution of dislocation 
 cell structures . for 
 partial , perfect , and shuffle 
 sets , screw dislocations always have the lowest energies . dissociated 
 glide dislocations always have lower energies than perfect glide dislocations . 

however , perfect glide dislocations do not necessarily dissociate 
 due to the energy barrier of dissociation .
in particular , the splitting 
 distance of perfect glide dislocations is seen to remain very small 
 in our dynamic md simulations .
the acute angle of the second lowest 
 energy ( shuffle , glide partial , glide perfect ) dislocations all equal 
 30 ( e.g. ,  = 30 , 150 , 330 ) .
screw dislocations are 
 much less mobile 
 than  edge dislocations . as a result ,
the screw dislocation 
 sectors perpendicular to the  sector are elongated when an 
  sector of a dislocation loop moves .
this suggests that screw 
 dislocations are likely to have a higher density than edge dislocations , 
 in good agreement with the low screw dislocation energies . without calculating the 
 energy profile 
 explicitly ,
we show that md simulations of dislocation motion at a 
 variety of stresses and temperatures can be used to effectively deduce 
 the activation free energy barrier and activation volume of thermally 
 activated glide of dislocations .
in particular , we found activation 
 free energy barriers of 0.14 and 0.27 ev and activation volumes of 
 17 and 36  , respectively , for  and  
 edge types of glide dislocations .
 cd1-xznxte 
 ( czt ) crystals are the leading semiconductors for radiation 
 detection , but their application is limited by the high cost of detector - grade 
 materials . 
 high crystal costs primarily result from property nonuniformity 
 that causes low manufacturing yield . 
 although tremendous efforts have 
 been made in the past to reduce te inclusions / precipitates in czt , 
 this has not resulted in an anticipated improvement in material property 
 uniformity . 
 moreover , it is recognized that in addition to te particles , 
 dislocation cells can also cause electric field perturbations and 
 the associated property nonuniformities . 
 further improvement of the 
 material , therefore , requires that dislocations in czt crystals be 
 understood and controlled . 
 here , we use a recently developed czt bond 
 order potential to perform representative molecular dynamics simulations 
 to study configurations , energies , and mobilities of 29 different 
 types of possible dislocations in cdte ( i.e. , x = 
 1 ) crystals . an efficient method to derive activation free energies 
 and activation volumes of thermally activated dislocation motion will 
 be explored . 
 our focus gives insight into understanding important 
 dislocations in the material and gives guidance toward experimental 
 efforts for improving dislocation network structures in czt crystals . 
Introduction
Interatomic Potential
Definition of Dislocation Types
Stacking Fault Energy
Dislocation Line Energies
Dislocation Loop Evolution under Shear
Dislocation Mobilities
Conclusion