Critical  review  on  “Grain-­‐boundary  engineering  markedly  reduces  susceptibility  to  intergranular  hydrogen  embrittlement  in  metallic  materials.”  

 Guoqiang  Xu  

 Department  of  Materials  Science  and  Engineering,  Massachusetts  Institute  of  

Technology  (MIT),  Cambridge,  MA  02139,  USA    S.   Bechtle,  M.   Kumar,   B.   P.   Somerday,  M.   E.   Launey,   R.   O.   Ritchie,   Grain-­‐boundary  engineering   markedly   reduces   susceptibility   to   intergranular   hydrogen  embrittlement  in  metallic  materials.    Acta  Mater.  57,  (2009)  4148-­‐4157  [1]    1.  Brief  Summary    

In   this   paper,   S.   Bechtle   et.   al.   investigated   the   susceptibility   of     a   nickel-­‐based   alloy   to   intergranular   hydrogen   embrittlement   (HE).   The   nickel   alloys   are  processed   using   cold   rolling   technique   with   different   controlling   parameters,  resulting   in   different   grain   boundary   (GB)   microstructures.   For   the   sample   with  high   fraction   of   ‘special’   GBs,   the   tensile   ductility   in   the   presence   of   hydrogen   is  almost   doubled   compared   to   that   in   samples   with   low   fraction   of   special   GBs.  Besides,   the   proportion   of   intergranular   fracture   was   significantly   lower   for   the  former.   The   author   attributes   the   reduction   in   the   severity   of   H-­‐induced  intergranular  embrittlement  to  the  fact  that  the  degree  of  hydrogen  segregation  at  special  GBs  is  lower.      2.  Introduction    

With   increasing  H   content,   the  mechanical   failure  mode   of  metals   changes,  going   from   ductile   transgranular,   to   ductile   intergranular,   to   brittle   intergranular  [2].   Each   of   these   transitions   is   associated   with   a   marked   reduction   of   overall  fracture   toughness   [3].   The   conditions   at   which   these   transitions   occur   and   the  degree   of   reduction   in   fracture   toughness   associated   with   them   are   functions   of  composition,   loading   configuration,   and   detailed   microstructure,   including   the  number  and   type  of   grain  boundaries  present   in   the  material   [4,  5].  GBs  exhibit   a  variety  of  structures  and  properties,  depending  on  their  crystallography  and  atomic-­‐level  state.  They  may  therefore  be  expected  to  exhibit  differing  susceptibilities  to  H-­‐induced  fracture.  GB  engineering  can  be  used  to  change  the  number  and  type  of  GB  in  metallic  alloys  and  therefore  to  manufacture  alloys  with  better  performance  in  H-­‐rich  environments  [6,  7].      3.  Literature  and  concepts    

Failure  upon  loading  may  occur  through  different  fracture  modes  depending  on   the   manner   in   which   the   crack   propagates   through   a   material.   In   ductile  materials,   crack   propagation   is   accompanied   by   a   large   amount   of   plastic  deformation   [8],   while   in   brittle   materials   cracks   moves   with   little   or   no   plastic  deformation  [9].  In  polycrystalline  materials,  cracks  propagation  through  the  grains  is  called  transgranular  fracture.  On  the  other  hand,  crack  propagation  along  GBs  is  termed   intergranular   fracture.   Fig.   1A   illustrates   intergranular   and   transgranular  fracture.      

 Fig.  1  Schematic  illustration  of  different  fracture  modes.  (A)  Intergranular  and  

transgranular  fracture.  (B)  Cleavage  and  ductile  fracture.    

Ductile  materials  usually  have  high  fracture  toughness  owing  to  the  work  of  plastic   deformation   expended   during   crack   advance.   Fracture   of   most   ductile  materials   proceeds   by   the   nucleation,   growth   and   coalescence   of   voids   (Fig.   1B)  [10].  Void  nucleation  usually  takes  place  in  regions  of  high  local  plastic  deformation  ahead   of   the   crack   tip,   resulting   in   a   fracture   surface   typically   covered   with  hemispherical   or   hemi-­‐ellipsoidal   dimples   [11].   Depending   on   the   loading  configuration  and  material  properties,  void  coalescence  may  occur  after  a  significant  amount   of   void   growth,  which   leads   to   a   fracture   surface   covered  with   large   and  deep  dimples,  or  start  soon  after  the  nucleation  of  voids,  giving  rise  to  shallow  and  small  dimples  [8,  12].      

Some   materials   like   diamond   can   undergo   pure   brittle   fracture   with   no  discernable  plasticity  associated  with  this  process  (Fig.  1B).  Because  no  plastic  work  is  expensed,  the  fracture  toughness  of  such  materials  is  therefore  low.  In  addition  to  

pure  brittle   fracture  where   there   is  no  plasticity,   fracture  can  also  be  preceded  by  micro-­‐plasticity,  but  void  nucleation  does  not  occur  [13].  One  example  of  this  type  of  fracture  is  when  dislocations  emitted  from  a  crack  tip  only  shield  the  crack  tip  stress  field   and   do   not   blunt   a   major   portion   of   the   crack   front,   allowing   it   to   remain  atomically  sharp  and  to  advance  by  atomic  decohesion  [9].  Another  example  is  slip-­‐induced  cleavage,  which  occurs  in  the  plastic  zone  ahead  of  and  away  from  the  crack  tip  [14].    

Theoretically,   ductile   to   brittle   transition   is   attributed   to   the   competition  between  two  processes,  decohesion  and  dislocation  emission,  as   is  shown  in  Fig.  2  [15-­‐17].   If  decohesion  process   is  easier   to  happen,  crack  will  propagate  by  slitting  atom  bonds   in   front   of   it   and   therefore   the   fracture  mode   is   brittle.   On   the   other  hand,  if  dislocation  emission  is  easier  to  happen,  cracks  will  get  blunted  due  to  the  dislocation   nucleation   with   little   extension.   At   the   same   time,   the   energy   is  dissipated   through  dislocation   activities   and   the  high   stress   fields   at   the   crack   tip  are  relieved,  further  preventing  its  propagation.  At  lower  temperatures  and/or  high  loads,   dislocation   activities   also   lead   to   void   formation,   resulting   in   crack  propagation  by  void  coalescence.    

 Fig.  2  The  competition  between  decohesion  and  dislocation  emission  determines  the  

ductile  to  brittle  transition  behavior  of  metals.      

Based   on   this   criterion,   H   may   influence   ductile   to   brittle   transition   by  influencing   the   decohesion   process   as   well   as   dislocation   emission   process.   H  generally   reduces   the   maximum   cohesive   force   between   two   atomic   layers   and  therefore  makes  it  easier  for  decohesion.  Carter  calculated  the  ideal  fracture  energy  of   iron   and   aluminum   in   the   presence   of   varying   amounts   of   H   using   periodic  density   functional   theory   [18].   He   found   that   the   ideal   fracture   energy   decreases  almost   linearly   with   increasing   H   coverage,   dropping   by   ~45%   at   one-­‐half  monolayer   coverage   and   indicating   a   substantial   reduction   of   cohesion   in   the  presence  of  H,  as  is  shown  in  Fig.  3.      

 Fig.  3  Decrease  in  normalized  surface  energies  !(!)/  !(0)  for  H  covered  

Al  (111)  and  Fe  (110)  as  hydrogen  coverage  !  increases.  (Figure  copied  from  Fig.  3  of  Ref.  [18])  

 Oriani  also  proposed  Hydrogen  Induced  Decohesion  as  a  mechanism  for  HE  

[19-­‐22]   and   supported   this   experimentally   by   tensile   loading   of   pre-­‐cracked   AISI  4340   steel   in   a  H   atmosphere   [22].   These   experiments   showed   that   as   the  partial  pressure  of  H  increases,  the  critical  stress  intensity  at  the  crack  tip  needed  to  restart  crack   propagation   decreases,   indicating   that   H   decreases   the   maximum   cohesive  force  between  atomic  layers  and  embrittles  these  crystals,  as  is  shown  in  Fig.  4.      

 Fig.  4  Pressures  of  hydrogen  and  deuterium,  marked  by  a  short  dash,  at  which  the  crack  remained  stationary  prior  to  the  pressures  at  which  the  crack  propagated,  

marked  by  a  circle  for  H,  and  a  square  for  D,  at  various  stress-­‐intensity  factors,  K.  (Figure  copied  from  Fig.  2  of  Ref.  [22])  

 On  the  other  hand,  hydrogen  may  influence  the  dislocation  process  through  

its   interaction  with  dislocations.  By   considering   the   size  misfit   strain   energy,  Rice  calculated   the  energy  release  rate   for   the  emission  of  a   straight-­‐line  dislocation   in  the  presence  of  H  [23].  Theoretical  analysis  showed  that  the  role  of  H  is  to  pin  the  movement  of  dislocations,  making  dislocation  emission  harder.  This  pinning  effect  favors  brittle  fracture.  In  addition  to  the  size  effect,  hydrogen  may  also  interact  with  dislocations   by   increasing   dislocation   mobility   and   allows   highly   localized  deformation   [24,  25].  This   is   the  Hydrogen-­‐Enhanced  Local  Plasticity  proposed  by  Beachem   as   a   mechanism   for   HE.   The   HELP   model   suggested   that   the   enhanced  plasticity  would  result   in   localized  softening,  which   leads   to   failure  by  plastic   flow  localization,   in   contrast   to   the   usual   sense   of   embrittlement.   Some   in   situ  environmental   cell   transmission   electron   microscopy   (TEM)   measurements   have  been  interpreted  as  lending  support  to  this  mechanism  [26].    

Different  GBs  have  different  susceptibilities  to  intergranular  HE.  For  example,  Wang  worked  out  the  degree  of  susceptibility  of  GBs  to  bismuth  embrittlement,  as  shown   in   Fig.   5.   The   results   clearly   indicated   that  Σ11   GB   is   more   susceptible   to  bismuth   embrittlement   than  Σ  9   and  Σ41   [27].   All   the   factors   that   influence   the  energy   release   rates   for   cleavage   and   dislocation   emission   may   play   a   role   in  determining  GB  susceptibility   to  HE.  Those   include   the  binding  energy  of  H  at   the  trapping   sites   of   the   GBs,   the   reduction   of   ideal   interface   separation   energies   at  certain  H  concentration,  as  well  as  the  shielding  of  GB  stress  field  in  the  presence  of  H.      

 Fig.  5  Predictions  of  the  Rice-­‐Thomson  model  for  cleavage  response  versus  

dislocation  emission  at  interfacial  cracks.  (Figure  copied  from  Fig.  13  of  Ref.  [27])    

4.  Discussion  on  the  results    

The   author   shows   that   GB   engineering   can   be   used   to   control   the  microstructures.   For   samples   of   the   same   size,   two   different   thermomechnaical  processes  are  adopted.  For  one  sample,  three  cycles  of  cold  rolling  (5%  reduction  in  the  thickness  per  cycle)  followed  by  a  15  min  anneal  at  900  degrees  in  air  furnace  (water   quenched),   resulting   a   microstructure   with   a   high   length   fraction   of   75%  special  boundaries  (62%  number  fraction).  For  the  other  sample,  four  cycles  of  cold  rolling  (20%  reduction  in  thickness  per  cycle)   followed  by  a  15  min  anneal  at  700  degrees   in   an   air   furnace   (water   quenched)   with   a   final   1   hour   grain-­‐coarsening  anneal  at  900  degrees,  resulting  in  a  low  length  fraction  of  46%  special  boundaries  (35%  number   fraction).  Moreover,  all   the  special  GBs   in  the  second  sample  are  !3  twin  boundaries.  Fig.  6,  taken  from  this  paper,  shows  the  EBSD  orientation  mapping  of  the  microstructures,  with  !3!  special  boundaries  in  blue  and  !3  twin  boundaries  in  red.  

 

 Fig.  6.  EBSD  orientation  maps  of  the  grain  structures  in  Ni-­‐201  of  the  

microstructures  with  a  (A)  low  fraction  (46%  by  length)  and  (B)  high  fraction  (75%  by  length)  of  special  grain  boundaries.  Random  boundaries  are  depicted  in  black,  R3  twin  boundaries  in  red,  and  other  R3n  special  boundaries  in  blue.  Grain  size  and  texture  remain  essentially  unchanged.  (Figure  copied  from  Fig.  1  of  the  paper)  

 The   interesting   question   arisen   from   this   result   is   that   what   is   the  

mechanism  for  using  GB  engineering  to  change  the  microstructures.  One  proposed  mechanism  is  called  the  ‘!3  regeneration  mechanism’  [28].  The  central  point  of  the  mechanism   is   that   when   two  !3s   boundaries   meet,   a  !9  boundary   is   formed   to  

complete   the   triple   junction.   When   such   a  !9  boundary   encounters   another  !3,   a  new  !3  boundary  (importantly,  not  a  coherent  annealing  twin)  will  be  regenerated  according  to  !3! + !3!!! → !3.    

 This   process   has   three   requirements.   First,   a   threshold   level   of  !3s   in  

materials  before  the  thermomechnaical  processing  is  required.  This  is  because  if  the  fraction  of  !3  GBs  is  too  small,  chances  are  small  for  two  !3s  to  meet  each  other  and  react.  This  also  indicates  that  profuse  annealing  twinning  is  a  prerequisite  for  most  GB   engineering.   Secondly,   the   thermomechnaical   process   has   to   be   iterative.  Otherwise,  the  regeneration  described  above  could  only  occur  once  and  therefore  it  is  not  able  for  the  sample  to  be  abundant  in  !3!  special  boundaries.  Last,  there  has  to   be   a   driving   force   for   selective   grain   boundary   migration   such   that   special  boundaries,  mainly  incoherent  !3s  and  some  !9s  will  migrate  preferentially.  This  is  why  annealing  must  be  followed  after  each  cold  rolling.  

 In  nanocrystalline  materials,   the volume fraction of GB (50% for 5 nm grains,

30% for 10 nm grains) reaches very high value and therefore the interface structure can dictate the macro mechanical properties [29]. However, the  average  grain  size  in  both  samples  in  this  work  is  around  30  !",  indicating  that  the  volume  fraction  of  GBs  is  extremely  small.   It   is   interesting  why  GBs  could  play  such  an   important   role  here.  One   of   the   reasons   may   be   that   H   prefers   to   segregate   to   the   GBs   and   cause  intergranular   fracture.   Crack   propagation   paths   are   within   the   connected   GB  networks,  which  makes  GB  type  and  structure  important  in  improving  the  fracture  property  of  materials.  

 Another   important   finding   in   this   work   is   that   in   the   presence   of   H  

concentrations   between   1200   and   3400   appm,   the   high   special   fraction  microstructure   showed  almost  double   the   tensile  ductility;   also,   the  proportion  of  intergranular   fracture  was  significantly   lower.   It’s   still  not  clear  how  those  special  GBs  can  influence  the  susceptibility  of  HE  in  alloys.  According  to  the  literature,  the  role   of   H   in   the   fracture   process   is   quite   complicated.   Not   only   can   it   reduce   the  maximum  cohesive  force  along  the  interface,  hence  facilitating  decohesion  process,  but  it  can  also  shield  the  interactions  between  dislocations,  increase  the  mobility  of  dislocations,  finally  leading  to  the  highly  localized  deformation  bands.    

 

 Fig.  7  Binding  energy  of  H  at  different  GBs,  computed  from  atomistic  simulations.  

(Figure  from  Ref.  [30])    

The  author  attribute  the  reduction  in  the  severity  of  H-­‐induced  intergranular  embrittlement   to   the   fact   that   the  binding   energy  of  H   at   special  !3!  GBs   is   small  and  therefore,   less  H  is  segregated  at  there  boundaries.  However,   literature  shows  that  the  binding  energy  of  H  at  !3!  is  not  differ  much  from  that  of  other  GBs.  Baskes  showed  using  atomistic  simulations  that  Σ3  (112)  and  Σ11  (113)  GBs  in  nickel  have  a  maximum  trapping  energy  of  0.24eV  while  at  Σ9  (221)  GBs  it  is  0.28eV  [30].  Hickel  studied   interaction   of   H   interstitials   with   GBs   in   alpha-­‐iron   and   gamma-­‐iron   by  density   functional   theory  and  showed  that   the  trapping  energies   for  H  at  Σ3  (112)  and  Σ11   (113)  GBs  are  0.18eV  while   that   for  Σ5   (310)  GB   is  much   lower  and  only  0.05eV   [31].   Ferthermore,   even   at   the   same   H   coverage,   the   reduction   of   ideal  interface  separation  energy  for  different  GBs  in  the  presence  of  H  may  be  different.  Hickel  showed  that  for  Σ3  GBs  in  alpha  iron,  the  interface  separation  energy  reduced  by  33.1%  while  for  Σ5  GBs  it  only  decreased  by  11.2%  [31].  

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