Strengthened　by　the　L12　precipitates　as　such　Ni3Al,　the　CuNi14Al3　alloyis　utilized　to　manufacture　the　shroud　rings　which　areone　of　the　key　components　of　power　generators.　However,　after　the　sequent　processing　procedures　including　vacuum　melting　casting,　forging　and　stress　relief　annealing　heat　treatment,　the　poor　surface　quality,　scrap　removal　induced　materials　waste,　forging　cracks　and　cavities　were　detected　from　time　to　time　in　the　final　products.　To　uncover　the　mechanism　of　the　formation　of　these　macroscopic　defects,　the　microstructural　features　and　mechanical　properties　of　the　CuNi14Al3　alloy　after　each　crucial　processing　stage　were　studied　by　optical　microscope　(OM),　scanning　electron　microscope　(SEM),　transmission　electron　microscope　(TEM),　energy　dispersive　X-ray　spectroscope　(EDS),　and　micro-hardness　tester.　Owing　to　the　solidification　characteristics　of　the　CuNi14Al3　alloy,　including　the　solidification　sequence,　rate　and　shrinkage,　the　microstructure　of　the　ingots　featured　of　coarse　dendritic　structures,　which　was　over　10　mm　in　length,　with　severe　elemental　segregation　and　micro-pores.　Since　nickel　had　higher　melting　point　than　copper,　the　precipitationinduced　elements　like　nickel　and　aluminum　to　segregate　in　the　dendritic　cores　with　a　segregation　ratio　up　to　0.3　and　0.4,　respectively.　From　the　dendritic　cores　to　the　interdendritic　regions,　the　shape　of　the　precipitate　particles　changed　gradually　from　large　regular-shaped　cuboids　(with　the　edge　length　around　200　nm)　to　fine　irregular　spherical　shapes,　with　an　average　diameter　of　around　20　nm.　The　hardness　values,　obtained　using　the　Vickers　indenter,　at　the　dendrite　cores　were　higher　than　those　in　the　interdendritic　regions　by　approximately　13.1%.　To　investigate　the　tensile　facture　behavior　of　the　ingot,　the　45　ocone-shaped　fractures　revealed　the　brittleness　characteristics　of　the　alloy.　However,　at　micro　scale,　the　facture　surface　was　dimple-shaped　morphology　with　mass　of　pores,　and　the　precipitate　particles　could　be　commonly　observed　at　the　bottom　of　the　dimples,　indicating　that　the　stress　concentration　around　the　casting　pores　and　shrinkage　led　to　the　fracture.　After　the　sequent　forging,　the　coarse　dendrites　and　the　micro-pores　were　inherited.　Moreover,　at　the　first　step　of　the　forging,　the　plastic　deformation　could　not　activate　sufficient　dynamic　recrystallization,　leading　to　the　bimodal　grain　size　distribution,　in　which　the　continuous　small　crystal　grains　distributed　among　the　giant　grains　inherited　from　the　casting　organization.　The　bimodal　microstructure　led　to　the　heterogeneity　of　the　deformation　partitions　at　further　forging　steps,　which,　in　turn,　induced　the　cracks　to　occur　at　the　interfacial　areas　between　the　giant　grains　and　the　micro-pores.　There　was　still　no　necking　observed　at　the　tensile　facture　of　the　forged　specimens.　Compared　with　those　of　the　casting　ingots,　the　dimples　became　shallower,　which　was　an　evidence　of　even　more　brittleness.　A　number　of　pores　initiating　from　the　casting　cavities　and　shrinkage　could　still　be　found,　and　the　pores　were　stretched　by　the　large　plastic　deformation.　All　the　defects,　including　dendritic　structure,　bimodal　distributed　grain　size,　pores　and　cracks,　were　left　to　the　last　processing　step,　stress　relief　annealing　heat　treatment.　Unfortunately　none　of　them　could　be　eliminated　by　the　annealing.　The　low　temperature　stress　relief　annealing　treatment　could　not　activate　the　recrystallization　sufficiently　and　many　giant　grains　were　left　over.　Due　to　the　high　forging　temperature,　which　was　higher　than　the　solutionizingheat　treatment　temperature　of　the　CuNi14Al3　alloy,　the　cracks　were　oxidized　as　Al2O3,　and　the　oxide　particles　along　the　cracks　impeded　the　closure　during　forging.　The　densities　of　the　alloys　were　measured　to　be　8.188　g•cm-3　after　casting,　8.245　g•cm-3　after　forging,　and　8.284　g•cm-3　after　stress　relief　annealing,　respectively.　All　values　were　significantly　lower　than　the　theoretical　value　(8.500　g•cm-3),　revealing　that　the　shrinkage　cavities　always　existed　and　the　defects　could　not　be　eliminated　completely　by　the　subsequent　hot　processing.　Although　the　cracks　emerging　during　the　forging　led　to　the　failure　of　the　product,　the　hardness　and　yield　strength　of　the　crack-free　parts　still　met　the　requirements,　mainly　resulting　from　the　dispersed　precipitation　hardening.　From　the　dendritic　cores　to　the　interdendritic　regions　of　the　casting　ingots,　the　size　of　the　Ni3Al　precipitate　particles　varied　from　200　to　20　nm.　The　crystal　structure　of　the　precipitation　was　similar　to　the　copper　matrix　and　had　the　same　orientation.　After　the　forging　and　the　annealing,　the　size　of　the　precipitation　became　uniform　at　around　20-50　nm,　which　guaranteed　the　higher　tensile　strength　at　855　MPa.　The　evolution　of　the　precipitate　particles　proved　that　the　forging　and　annealing　processing　was　effective　to　achieve　the　required　microstructure　and　mechanical　properties,　although　it　could　not　eliminate　the　casting　defects.It　could　be　concluded　that,　the　formation　of　all　the　defects　was　attributed　to　the　solidification　characteristics　of　the　alloy,　including　solidification　rate,　sequence　and　thermal　shrinkage.　Thus,　the　approach　to　improve　the　processing　could　be　raised.　Firstly,　to　improve　the　alloy　composition　by　adding　the　grain　refiner　to　promote　the　heteromorphy　nucleation　and/or　the　modifier　to　hinder　the　grain　growth　and　the　precipitating.　Secondly,　to　optimize　the　casting　processing　by　the　sloping　plate　casting　and　the　pressure　casting　to　refine　the　grain　size　and　restrain　the　shrinkage.　The　last　but　not　the　least,　to　add　another　homogenization　annealing　to　eliminate　the　dendritic　segregation　and　achieve　better　homogenization.　©　Editorial　Office　of　Chinese　Journal　of　Rare　Metals.　All　right　reserved.