Boom, G.,

Bronsveld, P.M.,

Hosson, J.Th.M. De,

Huis in't Veld, A.J.,

English

Huis in't Veld, A.J.

Boom, G.

Bronsveld, P.M.

De Hosson, J.Th.M.

NARCIS 

Superlattice Intrinsic Stacking Faults in γ' Precipitates

Dissertations of the University of Groningen

University of Groningen Digital Archive

Scripta  METALLURGICA  Vol.  19,  pp.  1123-1128,  1985  Pergamon  Press  Ltd. 
Printed  in the  U.S.A.  All  rights  reserved 
Superlattice Intrinsic Stacking Faults in "~' Precipitates 
A..].  Huis  in't Veld~ (3.  Boom~ P. M. Bronsveld, 3. Th.  M. De Hosson 
[Department of  Applied Physics,  Materials Science Centre, University  of  Oroningen, 
Nijenborgh  18,  9747  AG  Oroningen, The  Netherlands. 
(Received  June  24,  1985) 
1.  Introduction 
As  a  result  of  the  advance in gas  turbine  engine-technology  ~new  superalloys have  found  useful 
application  because  of  their  creep  resistance  at  high  temperature.  The  material  properties  of  MA  6000 
and  other  superalloys  depend  largely  upon  strengthening  by  L1  ordered y'  precipitates  and  to  a  lesser 
extent  upon  solid  solution  hardening.  Moreover, through  the  adc~ition  of  yttrium  oxide  particles,the  high 
temperature  performance  of  MA  6000  is  superior  to  conventional  superalloys [1].  This  paper  reports  on 
the  microstructure  of  MA  6000  deformed  at  760  °C  and  79Q °C. 
Ordered  Ni3Al  deforms  by  the  glide  of  a  pair  of  unit  dislocations  (:  one  supertattice 
dislocation),  which  minimizes  the  amount  of  disorder  by  reducing  the  APE{.  It  is  well  known  that  each 
unit  dislocation  can  split  into  two  Shockley  partials,  a/6<112>,  forming  a  so-called  complex  stacking 
fault  (CSF).  If  the ratio of  APB  and  5ISF  (Superlattice  Intrinsic  Stacking  Fault)  energies  is  high  enough, a 
superlattice  dislocation  dissociates  into  two  superlattice  Shocktey  partials,  a/3<112>~  separated  by  an 
SISF  as  observed  in  Zr.AI  [2].  Now  the  question  arises  which  dislocations  deform  a  superalloy containing 
a  high  volume  fractior~  of  L12  ordered  precipitates.  Mobile  dislocations  may  pass  the  precipitates  by 
shearing,  Orowan  bowing  or  (cooperative)  climb.  These  mechanisms  have  been  the  subjects  of  analyses 
which  mainly  consider  the  deformation  by  (pairs  of)  unit  dislocations  [3][4][5].  The  resistance  to 
dislocation  motion  of  an  alloy  strengthened by  50-60  volume percent  precipitates  is  highly  influenced  by 
the  capability  of  unit  dislocations  to  dissociate  into  partial  dislocations [6].  Especially,groups  of  partial 
dislocations  may  create  a  variety  of  fault  configurations  in  y'  precipitates.  Subsequent  partial  disloca- 
tions  on  the same slip  plane lead to  annihilation or  extension of  faulted  regions. 
This  paper  reports  upon  TEM  observations  at  200  kV  of  stacking  faults  confined  to  the  y' 
precipitates  of  MA  6000 after  deformation at  760  °C  and  790 °C. 
2.  Material and Experimental Procedure 
Chemical  composition of  MA  6000 (wt%) 
Ni  Cr  W  Mo  Al  Ti  Ta  C  B  Zr  Y203 
69  15  /4.0  2,0  /4,5  2.5  2.0  0.05  0.01  0.15  1,1 
Elemental  analysis  using  energy  and  wavelength  dispersive  X-ray  techniques  proved  small 
additions of  Re  and  Cu.  MA  6000  contain  columnar  grains  with  a  length  of  about  10  mm  which  share  a 
<110>  axis +  20  °  [7].  The  transverse grain  size  varies between 0.1  and  1  ram.  After  etching~  some details 
are  visible  within  the  columnar grains  with  an  optical  microscope.  With  X-ray  diffraction  and  TEM  these 
details  have  been  identified  to  be  small  angle  (5  ° )  boundaries.  X-ray  diffraction  on  powder,  chemically 
extracted,  gives  evidence  for  the  presence  of  yttrium  aluminium  oxide  (YA]O 3)  and  titanium  tungsten 
carbide (Ti,W)C.  About  50  volume percent  of  the  alloy  is  occupied by y'  (Llz)  Ni3(Al~Ti)  precipitates.  The 
small  misfit  between the y  matrix  and  y'  precipitates  causes  a  cuboida]  shape  of  the  large  precipitates 
which  have  edge  lengths  of  200-300  nm.  Between  the  cuboids  a  small  volume  fraction  of  y'  spheres  is 
found  with  a  mean  diameter  of  25  nm.  As  a  consequence  of  the  shearable  y'  spheres,  widely  separated 
superlattice  dislocations  are  found  in  the matrix  between the cuboids  in  MA  6000  as received. 
Test  specimens  are  prepared  with  the  gage  length  parallel  to  the  longitudinal  grain  direction. 
The  specimens contain  only  about  three columnar grains which  considerably reduced  the  lifetime  in  creep 
1123 
0036-9748/85  $3.00  +  .00 
Copyright  (c)  1985  Pergamon  Press  Ltd. 1124  STACKING  FAULTS  IN  y~  PRECIPITATES  Vol.  19,  No.  9 
experiments. Fracture  strain  was  found  to  be  approximately  8%.  2%  plastically  deformed  specimens  have 
been  used  for  TEM  observati~on{.  In  creep  2%  plastic  deformation  was  reached  at  the  onset  of  steady 
state  creep.  Tensile  (&  =  10-'s -')  and  creep  (538  MPa)  experiments  have  been  performed  at  760  °C  and 
790  °C,  respectively. 
At  the  fracture  surface  of  some  test  specimens,groups  of  spheres  are  found  (mean  diameter  5 
prn).  The  elemental  composition  of  these  spheres  could  not  be  distinguished  from  the  alloy  composition. 
Therefore,it  is  likely  that  the  spheres  consist  of  a  mechanically  alloyed  mixture  of  the  original  powders. 
In  figure  1  such  a  fracture  surface  is  depicted.  This  specimen  could  not  resist  368  MPa  at  790  °C.  It  may 
be  concluded  that  the  mechanical  strength  of  the  material  is  considerably  decreased  by  these  spheres. 
3.  Deformation  of  y' .Strenqthened  Alloy 
Upon  tensile  and  creep  deformation  at  760  °C  and  790  °C,  respectively,  single  unit  dislocations 
have  been  found  in  the  matrix.  At  the  y/y'  interface  small  ribbons  of  stacking  faults  result  from  the 
dissociation  into  partial  dislocations.  Inside  y'  precipitates  superdislocations  are  rarely  found.  The  main 
feature  of  the  microstructure  after  deformation  is  the  occurrence  of  stacking  faults  inside  y'  precipi- 
tates.  The  faults  can  be  divided  in  two  groups  which  are  depicted  in  figure  2.  By  tilting  it  has  been 
established  that  the  faults  are  located  on  two  {111}  planes.  As  a  result  of  the  <110>  grain  orientation, 
two  slip  systems  are  easily  activated.  Initially,the  stacking  faults  are  confined  to  the  y'  particles  which 
may  be proved by  two  micrographs  of  the  same  faulted  precipitate,  the  first  with  a  matrix  reflection  and 
the  second  with  a  superIattice  reflection  [8].  As  deformation  proceeds: faults  extend  into  the  matrix 
between  the  precipitates.  In  figure  2  it  can  be  seen  that  further  deformation  produces  microtwins  i.e. 
more  than  two  parallel  layers  of  stacking  faults.  At  fracture  the  microtwins  reach  a  considerable  density 
[9].  It  has  been  confirmed  experimentally  that  the  twins  are  parallel  to  the  macroscopically  visible  slip 
lines.  In  addition,it  has  been  established  that  the  direction  of  the  slip  lines  coincides  with  the  direction 
of  the  transgranular  fracture  surface  in  both  creep  and  tensile  experiments  at  760  °C. 
3a.  Stacking  faults  inside y'  precipitates 
An  experimental  method  to  verify  whether  the  faults  inside  y'  precipitates  are  CSF's  or  SISF's 
is  to  determine  the  Burgers  vector  of  the  leading  dislocation  of  a  fault.  The  determination  of  the  Burgers 
vector  has  been  performed  with  the  invisibility  criteria  g.b  =  0,  +  1/3,  -2/3,  4/3.  In  all  cases  the 
deviation  from  the  Bragg  reflection  was  taken  positive  (s  >  0).  The  results  are  summarized  in  table  1. 
The  invisibility  for  g.b  =  +  4/3  is  ambiguous.  Nevertheless,  a  comparison  of  fig.  3a  (g.b  =  +4/3)  with  fig. 
3b  (g.b  =  -4/3)  reveals  a  clear  difference  in  contrast.  Care  was  taken  that  on  reversing g  the  deviation s 
kept  the  same  magnitude. 
TABLE  1 
g.b  and g.R  analysis  of  the  fault  configuration  shown  in  fig.  3a-3d 
(i  =  invisible,  v  =  visible) 
The  extinction  criteria  for  a  stacking  fault  are cc =  n~ 
g.b  Observation  (Z=21T  g.R  Observation  Reference 
g  b=a/31112]  R=a/31112] 
220  413  i  813 ~  v  Fig.  3a 
220  -413  v  -813 ~  v  Fig.  3b 
200  -213  i  -413 ~  v  Fig.  3c 
200  2/3  v  4/3 ~  v  Fig. 3d 
202  2  v  4  ~  i 
020  2/3  v  4/3 ~  v 
020  -2/3  i  -4/3  ~  v 
The  occurrence  of  an  a/61112]  partial  dislocation  is  excluded  because  of  the  (in)visibility  of  the  leading 
dislocation  of  the  fault  in  figures  3a,  b  and  d. 
From  table  1  it  may  be  concluded  that  an  a/31112.]  partial  dislocation  creates  the  fault  inside 
the  y'  precipitate.  By  observing  the  outermost  fringes,the  intrinsic  character  of  the  faults  has  been 
proven.  The  dislocation-fault  configuration  shown  in  figures  3a-d  is  just  an  example  from  a  large 
number  with  the  same  extinction  behaviour. Vol.  19,  No,  9  STACKING  FAULTS  IN y'  PRECIPITATES  1125 
Fig.  1.  Group  of  spheres  at  a  fracture  sur- 
face  of  Ma  6000. 
Fig.  2.  Deformation  induced  stacking  faults 
inside y'  precipitates. 
a 
d 
Fig.  3.  a/31112]  dislocation  in front  of  a  SISF  inside y'  g.b  1/2  4/3,  -4/3,  -2/3,  2/3  in  figures  3a,  b,  c, 
dr  respectively.  (see  table  1). 1126  STACKING  FAULTS  IN  y'  PRECIPITATES  Vol.  19,  No.  9 
3b.  Growth  mechanism 
In  the  previous section  the  faults  have  been  identified  to  be  of  superlattice  intrinsic  nature.  In 
this  part  the  growth  mechanism  of  these  faults  will  be  analyzed.  Along  the  fringes  of  some  faults  a 
dislocation  segment  can  be  resolved.  In  figures  4a  and  4c  such  a  segment  is  visible.  In  figures  4b  and  4c 
it  can  be  seen  that  this  fault  is  extended  over  the  whole  y'  precipitate.  The  segment  seemed  to  be  the 
remainder  after  electropolishing  of  a  dislocation  loop  at  the  y/y'  interface.  With  the  invisibility  criteria 
of  section  3a,the  Burgers  vector  was  determined  to  be  a/61121],  see  table  2. 
TABLE  2 
g.b  analysis  of  the  dislocation  segment 
shown  in  figures  b,a-c 
(i  =  invisible,  v  =  visible) 
g.b  Obser-  Reference 
g  b=a/61121]  vation 
022  1  v  Fig. 4a 
~o  1/3  i 
20o  -1/3  i 
l"ff  1/3  i 
502  -2/3  i 
111  -2/3  i  Fig.  4b 
l fl  2/3  v  Fig.  4c 
f~l  o  i 
~[ToT] 
) 
1 
Fig.  5.  Growth  mechanism  of  stacking  faults 
inside y' precipitates. 
Although  the  plane  of  the  faults  in  figure  4  has  been  determined  by  tilting,  the  exact  value  of  the 
displacement  is  unknown.  The  three  possible  values  of  R  are:  a/312.11],  a/31112]  and  a/31121].  The  third 
would  imply  that  a/21121] is  operative  which  is  not  the  case  because  a/2<112>  partial  dislocations  cause 
other  fault  configurations  [6].  The  presence  of  an  a/61121] loop  together  with  R=a/31211]  or  a/31112] leads 
to  the  following  dissociation  mechanism  of  a  unit  dislocation  [10]:  a/21101] +  a/3['~15] +  a/61125].  This 
particular  dislocation  reaction  is  confirmed  by  the  presence  of  unit  dislocations  which  were  apparently 
obstructed  at  the  y/y'  interfaces. 
As  depicted  (schematically)  in  figure  5,  the  matrix  is  deformed  by  a/21105]  whereas  the 
precipitate  is  faulted  by  a/3121T].  At  the  y/y'  interface  an  a/61125]  dislocation  loop  is  left  behind. 
Obviously,  this  mechanism  accounts  for  faults  lying  inside  the  precipitates  only  as  experimentally 
observed  in  the  first  stage  of  deformation. 
3c.  Stair  rod  dislocations  in y'  precipitates 
As  a  result  of  the  <110>  grain  orientation,two  slip  systems  are  easily  activated.  In  the  first 
stage  of  plastic  deformation  only  one slip  system  is  operating,but  at  the  onset  of  steady  state  creep  both 
systems  are  active.  The  occurrence  of  two  slip  systems  causes  a  typical  fault  configuration  as  will  be 
described  below. 
Some  stacking  faults  (less  than  1%)  inside  y'  precipitates  reveal  a  line  of  intersection  of  two 
{111}  slip  planes.  By  edge-on  tilting  it  was  found  that  part  A  of  the  fault  in  fig.  6  is  located  on  (111) 
along [211]  and  part  B  on (111)  along [121].  It  is  likely  that  the  a/3<112>  partial  dislocations,  which  have 
created  the  faults,  have  formed  a  stair  rod  dislocation  at  the  line  of  intersection  of  the  two  slip  planes. 
in  figure  7  the  stair  rod  dislocation  is  visible.  Unfortunately,it  was  not  possible  to  identify  the  Burgers 
vector  of  the  stair  rod  dislocation  with  the  usual  invisibility  criteria.  If  two  a/3<112>  dislocations,  one 
on (111)  and  the  other  on  (1IS)react  along [110],several  stair  rod  dislocations  may  be  formed.  Regarding 
the  self  energy,  a/3<110>  is  the  most  likely  one. 
Assuming  the  leading  a/3<112>  partial  dislocations  which  have created  the  faults  to  be  of  pure 
edge  nature,the  exact  Burgers  vector  may  be  deduced  from  the  direction  of  the  faults.  This  implies  a 
dislocation  reaction  according  to  a/31211]  +  a/3[1ET]  -*  a/311-~0].  [n  silicon  grown-in  faults  have  been 
observed  with  a  similar  stair  rod  dislocation  [11].  The  stair  rod  in  figure  6  is  one  side  of  the  triangle 
which  reveals  a  contrast  typical  for  two  faults,one  above  another, which  have  a  symmetrical  inclination Vol.  19,  No.  9  STACKING  FAULTS  IN  y'  PRECIPITATES  1127 
Fig.  4.  a/6[121]  dislocation  along  the  frin- 
ges  of  a  SISF.  Figures  a,  b  and  c 
are  indicated  in  table  2. 
Fig,  6.  Stacking  fault  inside  y'.  Part  A 
along  [211]  on  (111).  Part  B  along 
[1;[1]  on (111). 
Fig.  7.  A  stair  rod  dislocation  is  visible 
along  the  line  of  intersection  of  the 
two  fault planes. 1128  STACKING  FAULTS  IN y'  PRECIPITATES  Vol.  19,  No.  9 
to  the  foil  as  explained by  Amelinckx  [12].  The  fringes  are  modulated parallel  and  perpendicular  to  the 
stair  rod  dislocation. 
4.  Discussion  and Conclusion 
After  2%  tensile and creep and  tensile deformation at  760 °C  and  790  °C,  respectively, a/3<112>  partial 
dislocations  have  been  identified  to  deform  y'  precipitates.  No  evidence was  found  for  a/2<112>  partial 
dislocations.  However,  fault  configurations  typical  for  a/2<112>  partial  dislocations  are  observed  after 
the  first  LCF  experiments at 760 °C. Upon further  deformation,  SlSF's  expand in  the  matrix  which  means 
that  the  energy  of  an  extended  fault  is  lower  than  the  energy  of  faulted  precipitates  separated  by 
unfauited  matrix.  The  extension  of  faults  in  the  matrix  might  be  stimulated  by  the  annihilation  of 
interfacia]  Shockley dislocations. 
From  our  results  it  might  be  concluded  that upon deformation at  760 °C  and  790  °C,  a/3<112> 
partial  dislocations  are  mobile  in  spite  of  a  high  Peierls  stress.  At  room  temperature  deformation  till 
fracture  produces a  rather low density of  twins.  As  a  matter  of  fach this  density  is  comparable with  the 
density  of  twins  in  the  as  received  material.  Calculations  show  a/3<112>  partial  dislocations  may  have 
glissile  and  sessile  core  structures  [13].  Therefore,the  explanation  might  be  that  at  room  temperature 
a/3<112>  partial dislocations are sessile. 
The  microstructure  shows  that  superlattice  dislocations  coupled  together  by  APB  only  play  a 
minor  role  in  deformation  of  MA  6000  at  760  °C  and  790  °C.  Yet  a  few  superdislocations  inside  y' 
precipitates  are  found.  Measuring  the  distance  between  the  two  dislocations  is  a  suitable  method  to 
obtain  an  estimate  of  the  APB  energy.  Unfortunately,  the  reliability  of  the  estimate  is  influenced  by 
residual  stresses.  Nevertheless, the  antiphase boundary energy  on  {111}  of  y'  precipitates  is  calculated  to 
be  144 (+ 10%)  mJ/m 2. 
In  a  recent  paper  about  the  so-called  sawtooth  faults  in  Ni_AI  [14]  Veyssi~.rre  has  calculated 
that  an  a/3<112>  partial  dislocation  has  a  stable  position  ~  '  '  at  a  certain  distance  from  a  Shockley  partial 
dislocation provided y  .../y.__  <  0 1  In  MA  6000  it  imnlies y  <  14  mJ/m 2  A  spontaneous dissociation 
•  br  ~t~  •  •  r  .  Sl  F.  " 
of  a  <110>  dlsloca~on  into  two  a/3<112>  partial  dlsloclaScllons  has  not  been  observed.  Therefore, 
ySlqF/Y~oB< >  0.03,  i.e.  YS[SF.  ' >  4  mJ/m 2.  In  addition  Veyssi~rre  states  that  an  a/3<112>  may  recombine 
wll~h  a#6  112>  by  tnermany  assisted  steps.  Obviously,a  similar  recombination  of  an  a/3<112>  with 
a/6<112>  is  necessary  to  account  for  faults  in  y'  precipitates  by  the  given  dissociation  and  recombina- 
tion  of  a  unit  dislocation.  The  observed  sawtooth  in  Ni3AI  and  faults  in  y'  precipitates  may  be 
interpreted as a  result  of  the same tendency to  dissociate. 
Acknowledqement 
The  work  is  part  of  the  research  program  of  the  Foundation  for  Fundamental  Research  on  Matter 
(F.O.M.-Utrecht)  and  has  been  made  possible  by  financial  support  from  the  Netherlands Organization  for 
the Advancement of  Pure Research (Z.W.O.-The  Hague). 
References 
1.  R.F.  Decker, 4th Int.  Symposium on Superalioys, Sept.  1980. 
2.  L.M.  Howe and M.  H.  Rainville, Proc.  9th int.  Congr.  on Electr.  Microscopy,  Toronto,  1978,  Vol.  1, 
404 (1978). 
3.  B.  Reppich,  Acta  Met. 31), 87 (1982). 
4.  M.  Mc Lean,  Acta  Met. 33, 545 (1985). 
5.  R.K.  Ham,  Proc.  3rd  Bolton Landing  Conf.,  Sept.  1969  (Eds.  B.  H.  Kear,  C.  T.  Sims,  N.  S.  Stoloff 
and  W.  H.  Westbrook),  365 (1969). 
6.  J.M.  Oblak and  B. H.  Kear, Electron  Microscopy  and Structure  of  Materials (Eds.  G.  Thomas,  R.  M. 
Fulrath  and R.  M. Fisher),  Un.  Calif.  Press,  566  (1972). 
7.  T.E.  Howson, D.  A.  Mervyn and J.  K.  Tien,  Met.  Trans.  11A,  1609 (1980). 
8.  A.J.  Huis  in't  Veld, P.  M.  Bronsveld, G.  Boom, J.  Th.  M.  De  Hosson,  7th  ICSMA,  Montreal, Quebec 
(Canada), August  1985. 
9.  A.J.  Huis  in't  Veld, P.  M.  Bronsveld, (3.  Boom,  J.  Th.  M.  De  Hosson,  Proc.  8th  European Congress 
on Electron  Microscopy, Budapest, Hungary,  Aug.  13-18,  Vol.  2,  851  (198/4). 
10.  B.H.  Kear,  A. F.  (3iamei, J.  M.  Silcock  and R.  K.  Ham, Scripta  Met. 2,  287 (1968). 
11.  H. F611, C.  B.  Carter and  M.  Wilkens,  Phys.  Stat. Sol. (a) 58,  393 (1980). 
12.  S.  Amelinckx, Solid State Physics,  Supplement 6.  The  Direct  Observation of  Dislocations,  Academic 
Press,  New York  and London (1964). 
13.  M.  Yamaguchi,  V.  Paidar,  D.  P.  Pope  and  V.  Vitek,  Phil.  Mag.  A  45,  No.  5,  883  (1982). 
14.  P.  Veyssi~rre, J.  Douin  and P.  Beauchamp, Phil.  Mag. A  51, No.  3,  469  (1985). 
15.  V.C.  Nardone and  J.  K.  Tien,  Scripta  Met. 17, 467  (1983). 

De Hosson,

Electron Microscopy and Structure of Materials (Eds.

Solid State Physics, Supplement 6. The Direct Observation of Dislocations,

ARTS repository - University of Groningen

Proceedings - University of Groningen

University of Groningen

University of Groningen Research Database

  
 University of Groningen
Superlattice Intrinsic Stacking Faults in γ' Precipitates
Huis in't Veld, A.J.; Boom, G.; Bronsveld, Paulus; De Hosson, J.T.M.
Published in:
Scripta Metallurgica
DOI:
10.1016/0036-9748(85)90021-3
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Citation for published version (APA):
Huis in't Veld, A. J., Boom, G., Bronsveld, P. M., & De Hosson, J. T. M. (1985). Superlattice Intrinsic
Stacking Faults in γ' Precipitates. Scripta Metallurgica, 19(9), 1123-1128. DOI: 10.1016/0036-
9748(85)90021-3
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Scripta METALLURGICA Vol. 19, pp. 1123-1128, 1985 Pergamon Press Ltd. 
Printed in the U.S.A. All rights reserved 
Superlattice Intrinsic Stacking Faults in "~' Precipitates 
A. . ] .  Huis in 't  Veld~ (3. Boom~ P. M. Bronsveld, 3. Th. M. De Hosson 
[Department of Applied Physics, Materials Science Centre, University of Oroningen, 
Nijenborgh 18, 9747 AG Oroningen, The Netherlands. 
(Received June 24, 1985) 
1. Introduction 
As a result of the advance in gas turbine engine-technology ~new superalloys have found useful 
application because of their creep resistance at high temperature. The material  properties of MA 6000 
and other superalloys depend largely upon strengthening by L1 ordered y' precipitates and to a lesser 
extent upon solid solution hardening. Moreover, through the adc~ition of yttr ium oxide part ic les,the high 
temperature performance of MA 6000 is superior to conventional superalloys [1]. This paper reports on 
the microstructure of MA 6000 deformed at 760 °C and 79Q °C. 
Ordered Ni3Al deforms by the glide of a pair of unit dislocations (:  one supertattice 
dislocation), which minimizes the amount of disorder by reducing the APE{. It is well known that each 
unit dislocation can split into two Shockley partials, a/6<112>, forming a so-called complex stacking 
fault (CSF). If the rat io of APB and 5ISF (Superlattice Intrinsic Stacking Fault) energies is high enough, a 
superlattice dislocation dissociates into two superlattice Shocktey partials, a/3<112>~ separated by an 
SISF as observed in Zr .A I  [2]. Now the question arises which dislocations deform a superalloy containing 
a high volume fractior~ of L12 ordered precipitates. Mobile dislocations may pass the precipitates by 
shearing, Orowan bowing or (cooperative) climb. These mechanisms have been the subjects of analyses 
which mainly consider the deformation by (pairs of) unit dislocations [3][4][5]. The resistance to 
dislocation motion of an alloy strengthened by 50-60 volume percent precipitates is highly influenced by 
the capabil i ty of unit dislocations to dissociate into partial  dislocations [6]. Especially,groups of partial  
dislocations may create a variety of fault configurations in y' precipitates. Subsequent partial  disloca- 
tions on the same slip plane lead to annihilation or extension of faulted regions. 
This paper reports upon TEM observations at 200 kV of stacking faults confined to the y' 
precipitates of MA 6000 after deformation at 760 °C and 790 °C. 
2. Material  and Experimental Procedure 
Chemical composition of MA 6000 (wt%) 
Ni Cr W Mo Al Ti Ta C B Zr Y203 
69 15 /4.0 2,0 /4,5 2.5 2.0 0.05 0.01 0.15 1,1 
Elemental analysis using energy and wavelength dispersive X-ray techniques proved small 
additions of Re and Cu. MA 6000 contain columnar grains with a length of about 10 mm which share a 
<110> axis + 20 ° [7]. The transverse grain size varies between 0.1 and 1 ram. After  etching~ some details 
are visible within the columnar grains with an optical microscope. With X-ray dif fract ion and TEM these 
details have been identif ied to be small angle (5 ° ) boundaries. X-ray dif fract ion on powder, chemically 
extracted, gives evidence for the presence of yttr ium aluminium oxide (YA]O 3) and t itanium tungsten 
carbide (Ti,W)C. About 50 volume percent of the alloy is occupied by y' (L lz)  Ni3(Al~Ti) precipitates. The 
small misf it  between the y matr ix and y' precipitates causes a cuboida] shape of the large precipitates 
which have edge lengths of 200-300 nm. Between the cuboids a small volume fraction of y' spheres is 
found with a mean diameter of 25 nm. As a consequence of the shearable y' spheres, widely separated 
superlattice dislocations are found in the matr ix between the cuboids in MA 6000 as received. 
Test specimens are prepared with the gage length parallel to the longitudinal grain direction. 
The specimens contain only about three columnar grains which considerably reduced the l i fet ime in creep 
1123 
0036-9748/85 $3.00 + .00 
Copyright (c) 1985 Pergamon Press Ltd. 
1124 STACKING FAULTS IN y~ PRECIP ITATES Vo l .  19, No. 9 
experiments. Fracture strain was found to be approximately 8%. 2% plastical ly deformed specimens have 
been used for TEM observati~on{. In creep 2% plastic deformation was reached at the onset of steady 
state creep. Tensile (& = 10-'s - ' )  and creep (538 MPa) experiments have been performed at 760 °C and 
790 °C, respectively. 
At the fracture surface of some test specimens,groups of spheres are found (mean diameter 5 
prn). The elemental composition of these spheres could not be distinguished from the alloy composition. 
Therefore, i t  is l ikely that the spheres consist of a mechanically alloyed mixture of the original powders. 
In figure 1 such a fracture surface is depicted. This specimen could not resist 368 MPa at 790 °C. It may 
be concluded that the mechanical strength of the material  is considerably decreased by these spheres. 
3. Deformation of y' .Strenqthened Al loy 
Upon tensile and creep deformation at 760 °C and 790 °C, respectively, single unit dislocations 
have been found in the matr ix.  At  the y/y '  interface small ribbons of stacking faults result from the 
dissociation into partial  dislocations. Inside y' precipitates superdislocations are rarely found. The main 
feature of the microstructure after deformation is the occurrence of stacking faults inside y' precipi-  
tates. The faults can be divided in two groups which are depicted in figure 2. By t i l t ing it has been 
established that the faults are located on two {111} planes. As a result of the <110> grain or ientat ion,  
two slip systems are easily activated. In i t ia l ly , the stacking faults are confined to the y' particles which 
may be proved by two micrographs of the same faulted precipitate,  the f irst with a matr ix ref lect ion and 
the second with a superIattice ref lect ion [8]. As deformation proceeds: faults extend into the matr ix 
between the precipitates. In figure 2 it can be seen that further deformation produces microtwins i.e. 
more than two parallel layers of stacking faults. At fracture the microtwins reach a considerable density 
[9]. It  has been confirmed experimental ly that the twins are parallel to the macroscopical ly visible slip 
lines. In addit ion, i t  has been established that the direction of the slip lines coincides with the direction 
of the transgranular fracture surface in both creep and tensile experiments at 760 °C. 
3a. Stacking faults inside y' precipitates 
An experimental method to veri fy whether the faults inside y' precipitates are CSF's or SISF's 
is to determine the Burgers vector of the leading dislocation of a fault. The determination of the Burgers 
vector has been performed with the invis ibi l i ty cr i ter ia g.b = 0, + 1/3, -2/3, 4/3. In all cases the 
deviation from the Bragg ref lect ion was taken positive (s > 0). The results are summarized in table 1. 
The invisibi l i ty for g.b = + 4/3 is ambiguous. Nevertheless, a comparison of fig. 3a (g.b = +4/3) with fig. 
3b (g.b = -4/3) reveals a clear difference in contrast. Care was taken that on reversing g the deviation s 
kept the same magnitude. 
TABLE 1 
g.b and g.R analysis of the fault configuration shown in fig. 3a-3d 
(i = invisible, v = visible) 
The ext inct ion cr i ter ia for a stacking fault are cc = n~ 
g.b Observation (Z=21T g.R Observation Reference 
g b=a/31112] R=a/31112] 
220 413 i 813 ~ v Fig. 3a 
220 -413 v -813 ~ v Fig. 3b 
200 -213 i -413 ~ v Fig. 3c 
200 2/3 v 4/3 ~ v Fig. 3d 
202 2 v 4 ~ i 
020 2/3 v 4/3 ~ v 
020 -2/3 i -4/3 ~ v 
The occurrence of an a/61112] part ial  dislocation is excluded because of the ( in)visibi l i ty of the leading 
dislocation of the fault  in figures 3a, b and d. 
From table 1 it  may be concluded that an a/31112.] part ial  dislocation creates the fault inside 
the y' precipitate.  By observing the outermost f r inges,the intrinsic character of the faults has been 
proven. The dislocation-fault  configuration shown in figures 3a-d is just an example from a large 
number with the same ext inct ion behaviour. 
Vol. 19, No, 9 STACKING FAULTS IN y' PRECIPITATES 1125 
Fig. 1. Group of spheres at a fracture sur- 
face of Ma 6000. 
Fig. 2. Deformation induced stacking faults 
inside y' precipitates. 
a 
d 
Fig. 3. a/31112] dislocation in front of a SISF inside y' g.b 1/2 4/3, -4/3, -2/3, 2/3 in figures 3a, b, c, 
dr respectively. (see table 1). 
1126 STACKING FAULTS IN y' PRECIP ITATES Vol. 19, No. 9 
3b. Growth mechanism 
In the previous section the faults have been ident i f ied to be of superlatt ice intr insic nature. In 
this part  the growth mechanism of these faults wi l l  be analyzed. Along the fringes of some faults a 
dis location segment can be resolved. In f igures 4a and 4c such a segment is visible. In figures 4b and 4c 
it  can be seen that this fault  is extended over the whole y' prec ip i tate.  The segment seemed to be the 
remainder a f ter  electropol ishing of a dislocation loop at the y /y '  interface.  With the inv is ib i l i ty  cr i ter ia  
of section 3a,the Burgers vector  was determined to be a/61121], see table 2. 
TABLE 2 
g.b analysis of the dislocation segment 
shown in figures b,a-c 
(i = invisible, v = visible) 
g.b Obser- Reference 
g b=a/61121] vation 
022 1 v Fig. 4a 
~o 1/3 i 
20o -1/3 i 
l " f f  1/3 i 
502 -2/3 i 
111 -2/3 i Fig. 4b 
l f l  2/3 v Fig. 4c 
f~l o i 
~[ToT] 
) 
1 
Fig. 5. Growth mechanism of stacking faults 
inside y' precipitates.  
Although the plane of the faults in f igure 4 has been determined by t i l t ing,  the exact  value of the 
displacement is unknown. The three possible values of R are: a/312.11], a/31112] and a/31121]. The third 
would imply that  a/21121] is operat ive which is not the case because a/2<112> part ia l  dislocations cause 
other fault  conf igurat ions [6]. The presence of an a/61121] loop together  wi th  R=a/31211] or a/31112] leads 
to the fo l lowing dissociation mechanism of a unit dis location [10]: a/21101] + a/3['~15] + a/61125]. This 
part icu lar  dislocation react ion is conf i rmed by the presence of unit dislocations which were apparent ly 
obstructed at the y /y '  interfaces. 
As depicted (schematical ly)  in f igure 5, the matr ix  is deformed by a/21105] whereas the 
prec ip i tate  is faulted by a/3121T]. At  the y /y '  inter face an a/61125] dis location loop is lef t  behind. 
Obviously, this mechanism accounts for faults lying inside the precip i tates only as exper imenta l ly  
observed in the f i rst  stage of deformat ion.  
3c. Stair rod dislocations in y' precip i tates 
As a result of the <110> grain o r ientat ion , two  slip systems are easily act ivated.  In the f i rst  
stage of plastic deformat ion only one slip system is operat ing,but  at the onset of steady state creep both 
systems are act ive.  The occurrence of two slip systems causes a typical  fault  conf igurat ion as wi l l  be 
described below. 
Some stacking faults (less than 1%) inside y' prec ip i tates reveal  a l ine of intersect ion of two 
{111} slip planes. By edge-on t i l t ing it  was found that part  A of the fault  in f ig. 6 is located on (111) 
along [211] and part  B on (111) along [121]. It is l ikely that the a/3<112> part ia l  dislocations, which have 
created the faults, have formed a stair rod dislocation at the l ine of intersect ion of the two slip planes. 
in f igure 7 the stair rod dislocation is visible. Unfor tunate ly , i t  was not possible to ident i fy  the Burgers 
vector  of the stair rod dislocation with the usual inv is ib i l i ty  cr i ter ia .  If  two a/3<112> dislocations, one 
on (111) and the other on (1IS)react along [110],several  stair rod dislocations may be formed. Regarding 
the self energy, a/3<110> is the most l ikely one. 
Assuming the leading a/3<112> part ia l  dislocations which have created the faults to be of pure 
edge nature , the exact Burgers vector  may be deduced f rom the direct ion of the faults. This implies a 
dis location react ion according to a/31211] + a/3[1ET] -* a/311-~0]. [n sil icon grown-in faults have been 
observed with a s imi lar  stair rod dislocation [11]. The stair rod in f igure 6 is one side of the tr iangle 
which reveals a contrast typical  for two faul ts ,one above another, which have a symmetr ica l  incl inat ion 
Vol. 19, No. 9 STACKING FAULTS IN y' PRECIP ITATES 1127 
Fig. 4. a/6[121] dis location along the fr in-  
ges of a SISF. Figures a, b and c 
are indicated in table 2. 
Fig, 6. Stack ing  fau l t  ins ide y'. Par t  A 
along [211] on (111). Part  B along 
[1;[1] on (111). 
Fig. 7. A stair rod dislocation is visible 
along the line of intersect ion of the 
two fault planes. 
1128 STACKING FAULTS IN y' PRECIPITATES Vol. 19, No. 9 
to the foil as explained by Amelinckx [12]. The fringes are modulated parallel and perpendicular to the 
stair rod dislocation. 
4. Discussion and Conclusion 
After  2% tensile and creep and tensile deformation at 760 °C and 790 °C, respectively, a/3<112> partial  
dislocations have been identif ied to deform y' precipitates. No evidence was found for a/2<112> partial  
dislocations. However, fault configurations typical for a/2<112> partial  dislocations are observed after 
the first LCF experiments at 760 °C. Upon further deformation, SlSF's expand in the matr ix which means 
that the energy of an extended fault is lower than the energy of faulted precipitates separated by 
unfauited matrix.  The extension of faults in the matr ix might be stimulated by the annihilation of 
interfacia] Shockley dislocations. 
From our results it might be concluded that upon deformation at 760 °C and 790 °C, a/3<112> 
partial  dislocations are mobile in spite of a high Peierls stress. At room temperature deformation t i l l  
fracture produces a rather low density of twins. As a matter of fach this density is comparable with the 
density of twins in the as received material .  Calculations show a/3<112> partial  dislocations may have 
glissile and sessile core structures [13]. Therefore,the explanation might be that at room temperature 
a/3<112> partial dislocations are sessile. 
The microstructure shows that superlattice dislocations coupled together by APB only play a 
minor role in deformation of MA 6000 at 760 °C and 790 °C. Yet  a few superdislocations inside y' 
precipitates are found. Measuring the distance between the two dislocations is a suitable method to 
obtain an estimate of the APB energy. Unfortunately, the rel iabi l i ty of the estimate is influenced by 
residual stresses. Nevertheless, the antiphase boundary energy on {111} of y' precipitates is calculated to 
be 144 (+ 10%) mJ/m 2. 
In a recent paper about the so-called sawtooth faults in Ni_AI [14] Veyssi~.rre has calculated 
that an a/3<112> part ial  dislocation has a stable position ~ ' ' at a certain distance from a Shockley partial  
dislocation provided y . . . / y . __  < 0 1 In MA 6000 it  imnlies y < 14 mJ/m 2 A spontaneous dissociation 
• br  ~ t ~  • • r . S l  F.  " 
of a <110> dlsloca~on into two a/3<112> partial  dlsloclaScllons has not been observed. Therefore, 
ySlqF/Y~oB< > 0.03, i.e. YS[SF. ' > 4 mJ/m 2. In addition Veyssi~rre states that an a/3<112> may recombine 
wll~h a#6 112> by tnermany assisted steps. Obviously,a similar recombination of an a/3<112> with 
a/6<112> is necessary to account for faults in y' precipitates by the given dissociation and recombina- 
tion of a unit dislocation. The observed sawtooth in Ni3AI and faults in y' precipitates may be 
interpreted as a result of the same tendency to dissociate. 
Acknowledqement 
The work is part of the research program of the Foundation for Fundamental Research on Matter 
(F.O.M.-Utrecht) and has been made possible by financial support from the Netherlands Organization for 
the Advancement of Pure Research (Z.W.O.-The Hague). 
References 
1. R .F .  Decker, 4th Int. Symposium on Superalioys, Sept. 1980. 
2. L .M.  Howe and M. H. Rainvi l le, Proc. 9th int. Congr. on Electr. Microscopy, Toronto, 1978, Vol. 1, 
404 (1978). 
3. B. Reppich, Acta Met. 31), 87 (1982). 
4. M. Mc Lean, Acta Met. 33, 545 (1985). 
5. R .K .  Ham, Proc. 3rd Bolton Landing Conf., Sept. 1969 (Eds. B. H. Kear, C. T. Sims, N. S. Stoloff 
and W. H. Westbrook), 365 (1969). 
6. J .M .  Oblak and B. H. Kear, Electron Microscopy and Structure of Materials (Eds. G. Thomas, R. M. 
Fulrath and R. M. Fisher), Un. Cal i f .  Press, 566 (1972). 
7. T .E .  Howson, D. A. Mervyn and J. K. Tien, Met. Trans. 11A, 1609 (1980). 
8. A . J .  Huis in't Veld, P. M. Bronsveld, G. Boom, J. Th. M. De Hosson, 7th ICSMA, Montreal, Quebec 
(Canada), August 1985. 
9. A . J .  Huis in't Veld, P. M. Bronsveld, (3. Boom, J. Th. M. De Hosson, Proc. 8th European Congress 
on Electron Microscopy, Budapest, Hungary, Aug. 13-18, Vol. 2, 851 (198/4). 
10. B .H .  Kear, A. F. (3iamei, J. M. Silcock and R. K. Ham, Scripta Met. 2, 287 (1968). 
11. H. F611, C. B. Carter and M. Wilkens, Phys. Stat. Sol. (a) 58, 393 (1980). 
12. S. Amelinckx, Solid State Physics, Supplement 6. The Direct Observation of Dislocations, Academic 
Press, New York and London (1964). 
13. M. Yamaguchi, V. Paidar, D. P. Pope and V. Vitek, Phil. Mag. A 45, No. 5, 883 (1982). 
14. P. Veyssi~rre, J. Douin and P. Beauchamp, Phil. Mag. A 51, No. 3, 469 (1985). 
15. V .C .  Nardone and J. K. Tien, Scripta Met. 17, 467 (1983). 


   University of GroningenSuperlattice Intrinsic Stacking Faults in γ' PrecipitatesHuis in't Veld, A.J.; Boom, G.; Bronsveld, P.M.; De Hosson, J.Th.M.Published in:Scripta MetallurgicaDOI:10.1016/0036-9748(85)90021-3IMPORTANT NOTE: You are advised to consult the publisher's version (publisher's PDF) if you wish to cite fromit. Please check the document version below.Document VersionPublisher's PDF, also known as Version of recordPublication date:1985Link to publication in University of Groningen/UMCG research databaseCitation for published version (APA):Huis in't Veld, A. J., Boom, G., Bronsveld, P. M., & De Hosson, J. T. M. (1985). Superlattice IntrinsicStacking Faults in γ' Precipitates. Scripta Metallurgica, 19(9), 1123-1128. https://doi.org/10.1016/0036-9748(85)90021-3CopyrightOther than for strictly personal use, it is not permitted to download or to forward/distribute the text or part of it without the consent of theauthor(s) and/or copyright holder(s), unless the work is under an open content license (like Creative Commons).Take-down policyIf you believe that this document breaches copyright please contact us providing details, and we will remove access to the work immediatelyand investigate your claim.Downloaded from the University of Groningen/UMCG research database (Pure): http://www.rug.nl/research/portal. For technical reasons thenumber of authors shown on this cover page is limited to 10 maximum.Download date: 12-11-2019Scripta METALLURGICA Vol. 19, pp. 1123-1128, 1985 Pergamon Press Ltd. Printed in the U.S.A. All rights reserved Superlattice Intrinsic Stacking Faults in "~' Precipitates A. . ] .  Huis in 't  Veld~ (3. Boom~ P. M. Bronsveld, 3. Th. M. De Hosson [Department of Applied Physics, Materials Science Centre, University of Oroningen, Nijenborgh 18, 9747 AG Oroningen, The Netherlands. (Received June 24, 1985) 1. Introduction As a result of the advance in gas turbine engine-technology ~new superalloys have found useful application because of their creep resistance at high temperature. The material  properties of MA 6000 and other superalloys depend largely upon strengthening by L1 ordered y' precipitates and to a lesser extent upon solid solution hardening. Moreover, through the adc~ition of yttr ium oxide part ic les,the high temperature performance of MA 6000 is superior to conventional superalloys [1]. This paper reports on the microstructure of MA 6000 deformed at 760 °C and 79Q °C. Ordered Ni3Al deforms by the glide of a pair of unit dislocations (:  one supertattice dislocation), which minimizes the amount of disorder by reducing the APE{. It is well known that each unit dislocation can split into two Shockley partials, a/6<112>, forming a so-called complex stacking fault (CSF). If the rat io of APB and 5ISF (Superlattice Intrinsic Stacking Fault) energies is high enough, a superlattice dislocation dissociates into two superlattice Shocktey partials, a/3<112>~ separated by an SISF as observed in Zr .A I  [2]. Now the question arises which dislocations deform a superalloy containing a high volume fractior~ of L12 ordered precipitates. Mobile dislocations may pass the precipitates by shearing, Orowan bowing or (cooperative) climb. These mechanisms have been the subjects of analyses which mainly consider the deformation by (pairs of) unit dislocations [3][4][5]. The resistance to dislocation motion of an alloy strengthened by 50-60 volume percent precipitates is highly influenced by the capabil i ty of unit dislocations to dissociate into partial  dislocations [6]. Especially,groups of partial  dislocations may create a variety of fault configurations in y' precipitates. Subsequent partial  disloca- tions on the same slip plane lead to annihilation or extension of faulted regions. This paper reports upon TEM observations at 200 kV of stacking faults confined to the y' precipitates of MA 6000 after deformation at 760 °C and 790 °C. 2. Material  and Experimental Procedure Chemical composition of MA 6000 (wt%) Ni Cr W Mo Al Ti Ta C B Zr Y203 69 15 /4.0 2,0 /4,5 2.5 2.0 0.05 0.01 0.15 1,1 Elemental analysis using energy and wavelength dispersive X-ray techniques proved small additions of Re and Cu. MA 6000 contain columnar grains with a length of about 10 mm which share a <110> axis + 20 ° [7]. The transverse grain size varies between 0.1 and 1 ram. After  etching~ some details are visible within the columnar grains with an optical microscope. With X-ray dif fract ion and TEM these details have been identif ied to be small angle (5 ° ) boundaries. X-ray dif fract ion on powder, chemically extracted, gives evidence for the presence of yttr ium aluminium oxide (YA]O 3) and t itanium tungsten carbide (Ti,W)C. About 50 volume percent of the alloy is occupied by y' (L lz)  Ni3(Al~Ti) precipitates. The small misf it  between the y matr ix and y' precipitates causes a cuboida] shape of the large precipitates which have edge lengths of 200-300 nm. Between the cuboids a small volume fraction of y' spheres is found with a mean diameter of 25 nm. As a consequence of the shearable y' spheres, widely separated superlattice dislocations are found in the matr ix between the cuboids in MA 6000 as received. Test specimens are prepared with the gage length parallel to the longitudinal grain direction. The specimens contain only about three columnar grains which considerably reduced the l i fet ime in creep 1123 0036-9748/85 $3.00 + .00 Copyright (c) 1985 Pergamon Press Ltd. 1124 STACKING FAULTS IN y~ PRECIP ITATES Vo l .  19, No. 9 experiments. Fracture strain was found to be approximately 8%. 2% plastical ly deformed specimens have been used for TEM observati~on{. In creep 2% plastic deformation was reached at the onset of steady state creep. Tensile (& = 10-'s - ' )  and creep (538 MPa) experiments have been performed at 760 °C and 790 °C, respectively. At the fracture surface of some test specimens,groups of spheres are found (mean diameter 5 prn). The elemental composition of these spheres could not be distinguished from the alloy composition. Therefore, i t  is l ikely that the spheres consist of a mechanically alloyed mixture of the original powders. In figure 1 such a fracture surface is depicted. This specimen could not resist 368 MPa at 790 °C. It may be concluded that the mechanical strength of the material  is considerably decreased by these spheres. 3. Deformation of y' .Strenqthened Al loy Upon tensile and creep deformation at 760 °C and 790 °C, respectively, single unit dislocations have been found in the matr ix.  At  the y/y '  interface small ribbons of stacking faults result from the dissociation into partial  dislocations. Inside y' precipitates superdislocations are rarely found. The main feature of the microstructure after deformation is the occurrence of stacking faults inside y' precipi-  tates. The faults can be divided in two groups which are depicted in figure 2. By t i l t ing it has been established that the faults are located on two {111} planes. As a result of the <110> grain or ientat ion,  two slip systems are easily activated. In i t ia l ly , the stacking faults are confined to the y' particles which may be proved by two micrographs of the same faulted precipitate,  the f irst with a matr ix ref lect ion and the second with a superIattice ref lect ion [8]. As deformation proceeds: faults extend into the matr ix between the precipitates. In figure 2 it can be seen that further deformation produces microtwins i.e. more than two parallel layers of stacking faults. At fracture the microtwins reach a considerable density [9]. It  has been confirmed experimental ly that the twins are parallel to the macroscopical ly visible slip lines. In addit ion, i t  has been established that the direction of the slip lines coincides with the direction of the transgranular fracture surface in both creep and tensile experiments at 760 °C. 3a. Stacking faults inside y' precipitates An experimental method to veri fy whether the faults inside y' precipitates are CSF's or SISF's is to determine the Burgers vector of the leading dislocation of a fault. The determination of the Burgers vector has been performed with the invis ibi l i ty cr i ter ia g.b = 0, + 1/3, -2/3, 4/3. In all cases the deviation from the Bragg ref lect ion was taken positive (s > 0). The results are summarized in table 1. The invisibi l i ty for g.b = + 4/3 is ambiguous. Nevertheless, a comparison of fig. 3a (g.b = +4/3) with fig. 3b (g.b = -4/3) reveals a clear difference in contrast. Care was taken that on reversing g the deviation s kept the same magnitude. TABLE 1 g.b and g.R analysis of the fault configuration shown in fig. 3a-3d (i = invisible, v = visible) The ext inct ion cr i ter ia for a stacking fault are cc = n~ g.b Observation (Z=21T g.R Observation Reference g b=a/31112] R=a/31112] 220 413 i 813 ~ v Fig. 3a 220 -413 v -813 ~ v Fig. 3b 200 -213 i -413 ~ v Fig. 3c 200 2/3 v 4/3 ~ v Fig. 3d 202 2 v 4 ~ i 020 2/3 v 4/3 ~ v 020 -2/3 i -4/3 ~ v The occurrence of an a/61112] part ial  dislocation is excluded because of the ( in)visibi l i ty of the leading dislocation of the fault  in figures 3a, b and d. From table 1 it  may be concluded that an a/31112.] part ial  dislocation creates the fault inside the y' precipitate.  By observing the outermost f r inges,the intrinsic character of the faults has been proven. The dislocation-fault  configuration shown in figures 3a-d is just an example from a large number with the same ext inct ion behaviour. Vol. 19, No, 9 STACKING FAULTS IN y' PRECIPITATES 1125 Fig. 1. Group of spheres at a fracture sur- face of Ma 6000. Fig. 2. Deformation induced stacking faults inside y' precipitates. a d Fig. 3. a/31112] dislocation in front of a SISF inside y' g.b 1/2 4/3, -4/3, -2/3, 2/3 in figures 3a, b, c, dr respectively. (see table 1). 1126 STACKING FAULTS IN y' PRECIP ITATES Vol. 19, No. 9 3b. Growth mechanism In the previous section the faults have been ident i f ied to be of superlatt ice intr insic nature. In this part  the growth mechanism of these faults wi l l  be analyzed. Along the fringes of some faults a dis location segment can be resolved. In f igures 4a and 4c such a segment is visible. In figures 4b and 4c it  can be seen that this fault  is extended over the whole y' prec ip i tate.  The segment seemed to be the remainder a f ter  electropol ishing of a dislocation loop at the y /y '  interface.  With the inv is ib i l i ty  cr i ter ia  of section 3a,the Burgers vector  was determined to be a/61121], see table 2. TABLE 2 g.b analysis of the dislocation segment shown in figures b,a-c (i = invisible, v = visible) g.b Obser- Reference g b=a/61121] vation 022 1 v Fig. 4a ~o 1/3 i 20o -1/3 i l " f f  1/3 i 502 -2/3 i 111 -2/3 i Fig. 4b l f l  2/3 v Fig. 4c f~l o i ~[ToT] ) 1 Fig. 5. Growth mechanism of stacking faults inside y' precipitates.  Although the plane of the faults in f igure 4 has been determined by t i l t ing,  the exact  value of the displacement is unknown. The three possible values of R are: a/312.11], a/31112] and a/31121]. The third would imply that  a/21121] is operat ive which is not the case because a/2<112> part ia l  dislocations cause other fault  conf igurat ions [6]. The presence of an a/61121] loop together  wi th  R=a/31211] or a/31112] leads to the fo l lowing dissociation mechanism of a unit dis location [10]: a/21101] + a/3['~15] + a/61125]. This part icu lar  dislocation react ion is conf i rmed by the presence of unit dislocations which were apparent ly obstructed at the y /y '  interfaces. As depicted (schematical ly)  in f igure 5, the matr ix  is deformed by a/21105] whereas the prec ip i tate  is faulted by a/3121T]. At  the y /y '  inter face an a/61125] dis location loop is lef t  behind. Obviously, this mechanism accounts for faults lying inside the precip i tates only as exper imenta l ly  observed in the f i rst  stage of deformat ion.  3c. Stair rod dislocations in y' precip i tates As a result of the <110> grain o r ientat ion , two  slip systems are easily act ivated.  In the f i rst  stage of plastic deformat ion only one slip system is operat ing,but  at the onset of steady state creep both systems are act ive.  The occurrence of two slip systems causes a typical  fault  conf igurat ion as wi l l  be described below. Some stacking faults (less than 1%) inside y' prec ip i tates reveal  a l ine of intersect ion of two {111} slip planes. By edge-on t i l t ing it  was found that part  A of the fault  in f ig. 6 is located on (111) along [211] and part  B on (111) along [121]. It is l ikely that the a/3<112> part ia l  dislocations, which have created the faults, have formed a stair rod dislocation at the l ine of intersect ion of the two slip planes. in f igure 7 the stair rod dislocation is visible. Unfor tunate ly , i t  was not possible to ident i fy  the Burgers vector  of the stair rod dislocation with the usual inv is ib i l i ty  cr i ter ia .  If  two a/3<112> dislocations, one on (111) and the other on (1IS)react along [110],several  stair rod dislocations may be formed. Regarding the self energy, a/3<110> is the most l ikely one. Assuming the leading a/3<112> part ia l  dislocations which have created the faults to be of pure edge nature , the exact Burgers vector  may be deduced f rom the direct ion of the faults. This implies a dis location react ion according to a/31211] + a/3[1ET] -* a/311-~0]. [n sil icon grown-in faults have been observed with a s imi lar  stair rod dislocation [11]. The stair rod in f igure 6 is one side of the tr iangle which reveals a contrast typical  for two faul ts ,one above another, which have a symmetr ica l  incl inat ion Vol. 19, No. 9 STACKING FAULTS IN y' PRECIP ITATES 1127 Fig. 4. a/6[121] dis location along the fr in-  ges of a SISF. Figures a, b and c are indicated in table 2. Fig, 6. Stack ing  fau l t  ins ide y'. Par t  A along [211] on (111). Part  B along [1;[1] on (111). Fig. 7. A stair rod dislocation is visible along the line of intersect ion of the two fault planes. 1128 STACKING FAULTS IN y' PRECIPITATES Vol. 19, No. 9 to the foil as explained by Amelinckx [12]. The fringes are modulated parallel and perpendicular to the stair rod dislocation. 4. Discussion and Conclusion After  2% tensile and creep and tensile deformation at 760 °C and 790 °C, respectively, a/3<112> partial  dislocations have been identif ied to deform y' precipitates. No evidence was found for a/2<112> partial  dislocations. However, fault configurations typical for a/2<112> partial  dislocations are observed after the first LCF experiments at 760 °C. Upon further deformation, SlSF's expand in the matr ix which means that the energy of an extended fault is lower than the energy of faulted precipitates separated by unfauited matrix.  The extension of faults in the matr ix might be stimulated by the annihilation of interfacia] Shockley dislocations. From our results it might be concluded that upon deformation at 760 °C and 790 °C, a/3<112> partial  dislocations are mobile in spite of a high Peierls stress. At room temperature deformation t i l l  fracture produces a rather low density of twins. As a matter of fach this density is comparable with the density of twins in the as received material .  Calculations show a/3<112> partial  dislocations may have glissile and sessile core structures [13]. Therefore,the explanation might be that at room temperature a/3<112> partial dislocations are sessile. The microstructure shows that superlattice dislocations coupled together by APB only play a minor role in deformation of MA 6000 at 760 °C and 790 °C. Yet  a few superdislocations inside y' precipitates are found. Measuring the distance between the two dislocations is a suitable method to obtain an estimate of the APB energy. Unfortunately, the rel iabi l i ty of the estimate is influenced by residual stresses. Nevertheless, the antiphase boundary energy on {111} of y' precipitates is calculated to be 144 (+ 10%) mJ/m 2. In a recent paper about the so-called sawtooth faults in Ni_AI [14] Veyssi~.rre has calculated that an a/3<112> part ial  dislocation has a stable position ~ ' ' at a certain distance from a Shockley partial  dislocation provided y . . . / y . __  < 0 1 In MA 6000 it  imnlies y < 14 mJ/m 2 A spontaneous dissociation • br  ~ t ~  • • r . S l  F.  " of a <110> dlsloca~on into two a/3<112> partial  dlsloclaScllons has not been observed. Therefore, ySlqF/Y~oB< > 0.03, i.e. YS[SF. ' > 4 mJ/m 2. In addition Veyssi~rre states that an a/3<112> may recombine wll~h a#6 112> by tnermany assisted steps. Obviously,a similar recombination of an a/3<112> with a/6<112> is necessary to account for faults in y' precipitates by the given dissociation and recombina- tion of a unit dislocation. The observed sawtooth in Ni3AI and faults in y' precipitates may be interpreted as a result of the same tendency to dissociate. Acknowledqement The work is part of the research program of the Foundation for Fundamental Research on Matter (F.O.M.-Utrecht) and has been made possible by financial support from the Netherlands Organization for the Advancement of Pure Research (Z.W.O.-The Hague). References 1. R .F .  Decker, 4th Int. Symposium on Superalioys, Sept. 1980. 2. L .M.  Howe and M. H. Rainvi l le, Proc. 9th int. Congr. on Electr. Microscopy, Toronto, 1978, Vol. 1, 404 (1978). 3. B. Reppich, Acta Met. 31), 87 (1982). 4. M. Mc Lean, Acta Met. 33, 545 (1985). 5. R .K .  Ham, Proc. 3rd Bolton Landing Conf., Sept. 1969 (Eds. B. H. Kear, C. T. Sims, N. S. Stoloff and W. H. Westbrook), 365 (1969). 6. J .M .  Oblak and B. H. Kear, Electron Microscopy and Structure of Materials (Eds. G. Thomas, R. M. Fulrath and R. M. Fisher), Un. Cal i f .  Press, 566 (1972). 7. T .E .  Howson, D. A. Mervyn and J. K. Tien, Met. Trans. 11A, 1609 (1980). 8. A . J .  Huis in't Veld, P. M. Bronsveld, G. Boom, J. Th. M. De Hosson, 7th ICSMA, Montreal, Quebec (Canada), August 1985. 9. A . J .  Huis in't Veld, P. M. Bronsveld, (3. Boom, J. Th. M. De Hosson, Proc. 8th European Congress on Electron Microscopy, Budapest, Hungary, Aug. 13-18, Vol. 2, 851 (198/4). 10. B .H .  Kear, A. F. (3iamei, J. M. Silcock and R. K. Ham, Scripta Met. 2, 287 (1968). 11. H. F611, C. B. Carter and M. Wilkens, Phys. Stat. Sol. (a) 58, 393 (1980). 12. S. Amelinckx, Solid State Physics, Supplement 6. The Direct Observation of Dislocations, Academic Press, New York and London (1964). 13. M. Yamaguchi, V. Paidar, D. P. Pope and V. Vitek, Phil. Mag. A 45, No. 5, 883 (1982). 14. P. Veyssi~rre, J. Douin and P. Beauchamp, Phil. Mag. A 51, No. 3, 469 (1985). 15. V .C .  Nardone and J. K. Tien, Scripta Met. 17, 467 (1983). 

Superlattice Intrinsic Stacking Faults in γ' Precipitates

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