Improved radiation resistance in metals via adaptive martensitic transformation

Domain: physics
Published: 2025-12-11
Editorial outcome: minor_revision

Abstract
Materialstypicallyexperienceseriousmicrostructureandperformance
degradationsunderirradiationinnuclearreactors.Toexploreradiation-
resistantmetalswithhighdesignflexibilityisurgentlyrequestedforthesafe
applicationofnuclearenergy.Inthiswork,wediscoverananti-radiation
mechanismforthispurposeinagradientnanostructurednucleargradeaus-
teniticstainlesssteelpreparedbyaflexiblesurfacenano-crystallization
approach.Aspecial3-dimensionalmicrostructurenetwork,consistingoflow-
energygrainboundaries,stackingfaults,anddislocationnetworks,isintro-
ducedinthenanostructure,sothatalarge-scaleadaptivemartensitictrans-
formationmechanismisactivatedunderirradiationevenatextremelyhigh
radiationdosesandhightemperatures.Consequently,theradiationresistance
issignificantlyenhanced,whileasuperiormechanicalpropertyisretained,in
nanostructuredsamplescomparedtocoarse-grainedcounterparts.Results
presentedinthisworkthusexploreastrategytoprepareradiation-resistant
metalsinfuture.

Introduction
Nuclearenergyhasreceivedworldwideattentioninpastdecadesasa
powerfulsourcetosolveenergyandenvironmentalproblemsonour
planet. To use nuclear energy in a secure manner, the structural
materialshavetobereliableduringtheirserviceinnuclearreactors.
Moreover,the requirementofextending the lifeofexisting nuclear
powerplantsmeansmuchlongerservicetimeofmaterialsunderthe
extreme radiation environment in future1–6. Therefore, structural
materials with enhanced radiation resistance are expected to be
developed, and researchers have committed to this issue in recent
years. For example, the role of grain boundaries (GBs) to increase
radiation resistance has been revealed by experiments and
simulations7–9.Inspiredbysuchworks,decreasingradiationdefectsby
a high density of GBs has been intensively investigated with the
development of nanocrystalline (NC) materials10. However, some
resultsinsistedthattheimprovementbyGBtrappingdefectsislimited
in NC materials, mostly due to the distinct graingrowthevenunder
low radiation doses11–13. In addition, trapping radiation defects by
the structural interfaces between nanosized particles and matrix has become a
promising method recently, represented by developing oxide
dispersion-strengthenedsteels14.Butthismethodislimitedbyitslow
productionefficiencyandlowbatchstability.Similarly,single-phase
concentrated solid-solution alloys (e.g., high-entropy alloys, HEAs)
havealsobeensuggestedtoexhibithighradiationresistance15,but
theirapplicationsarelimitedbyhighcostandlowstabilityathigh
temperaturesandhighdoses16.
It is noted that enhancing radiation resistance means not only
decreasingthedensityofradiationdefectsbutalsomaintainingthe
excellentserviceperformanceofmaterials.Thelocalmicrostructure
aroundinterfaceswouldbechangedbyabsorbingradiationdefects,
seriously degrading the mechanical properties of materials and
leadingintofailureduringserviceinreactors17–19.Therefore,adjusting
strategy by controlling both mechanical properties and radiation
defect density is crucial to the development of advanced nuclear
structuralmaterials.Forexample,aself-healingapproachwasintro-
ducedinaHEAwithtransformationinducedplasticity(TRIP)effectby
balancing radiation induced martensitic transformation with tem-
perature induced reverse transformation, so that a high stability of
mechanicalpropertieswasachievedinthematerialunderirradiation20.
Inthiswork,throughexperimentsandmoleculardynamics(MD)
simulations, we discovered that a gradient nanostructured (GNS)
austenitic stainless steel exhibits an extraordinary radiation resis-
tance, including decreased radiation defects and enhanced
mechanical properties, even under extremely high radiation doses
andhightemperatures.Moreover,anovelanti-radiationmechanism
throughalarge-scalephasetransformationhasbeenproposedand
confirmed.

Results
We used a commercial nuclear grade 304 austenitic stainless steel
(304NG),whichis widely used in current nuclearreactorsand con-
sidered as candidate material in future reactors, mostly due to the
outstandingperformanceofmechanicalproperties,high-temperature
resistance,andcorrosionresistance,andprocessedplatesamplesby
usingsurfacemechanicalrollingtreatment(SMRT)tointroduceaGNS
surfacelayer.BothGNSandcoarse-grained(CG)samplesunderwent
simultaneousironionradiationproceduresonasamesampleholder,
withthemaximumradiationdoseupto155displacementsperatom
(dpa). Accordingly, the ion implantation range is within 0−4.7μm
whilethepeakdamageisat~3.96μminirradiatedsamples.Detailscan
befoundin“Methods”,SupplementaryFigs.1–4andSupplementary
Tables1and2.Itisnotedthationimplantationexperimentsarehighly
valuableforunderstandingfundamentalaspectsofradiationeffectsas
the 1st step to explore radiation resistant materials under extreme
radiation environments, considering the cost, time and protection
neededagainstradiationinneutronirradiations21,22.
Transmissionelectronmicroscopy(TEM)observationsrevealthat
thesurfacelayertoendurethepeakradiationdamageconsistsoflath
structurewithanaveragetransversalsizeof37±13nminas-prepared
GNS sample (Fig. 1a, b). Through transmission Kikuchi diffraction
(TKD)measurements(Fig.1c),wefoundthatthedominantinterfaces
between laths are Σ3 (111) GBs and low-angle GBs (LAGBs). Further
characterizationsbyusingscanningTEM(STEM)indicatethatafew
intersectingstackingfaults(SFs,withadensityof~7.5×1014m−2)are
emittedfromtheGBsintothelaths(Fig.1dandSupplementaryFig.5).
Meanwhile,numerousdislocationnetworks(DNs,asmarkedinFig.1d)
existinthegrains.Thisfeatureisalsorevealedbya3-dimensional(3D)
model constructed at atomic scale by simulating SMRT in
304NG(Fig.1e).

Inmetalsunderirradiation,dislocationloops(DLs)areonekindofkey
defectstructuresinunderstandingtheresultedproperties(e.g.,creep,
swelling,hardening,andfatigue).Therefore,thedensityandsizeof
DLsaretypicallyusedtoestimatetheradiationdamagesinmaterials
underirradiationatlowtemperatures23,24.Inthiswork,ahighdensity
ofDLsformsinirradiatedCGsampleswhilenoremarkableDLsformin
irradiatedGNSsamplesevenatanextremelyhighdoseof155dpaat
roomtemperature,asshowninFig.2a,b,andSupplementaryFig.6.
Meanwhile, it isnoticed that the meangrain size is still in the nan-
ometer scale in irradiated GNS samples, and distinct martensitic
transformationoccursinbothgroupsofsamples(asreflectedbythe
increasing diffraction intensity of martensite in the corresponding
SAEDpatterns).
To quantify the evolution of defects under irradiation, all the
irradiatedsampleswerecarefullyanalyzedbycombiningTEMbright-
field (BF) and weak-beam dark-field (WBDF) observations (Supple-
mentaryFig.7).Subsequently,thecorrelationsofthedensityandsize
of DLs with radiation dose are summarized in Fig. 2c and Supple-
mentaryTable3.WecanseeboththemeandensityandsizeofDLs
(andthecorrespondinguncertainties)inirradiatedGNSsamplesare
typicallymuchlowerthanthoseinirradiatedCGcounterpartsatdoses
>5dpa. For example, the density is (5.96±0.53)×1020m−3 and the
diameteris4.69±0.78nmforDLsinGNSsampleat155dpa.Themean
valuesareonly~3.8%and~63%,respectively,ofthoseinCGcounter-
partafterthesameirradiation.Meanwhile,thedensityandsizeofDLs
significantly decrease with increasing dose within the early stage
(2–10dpa,seeSupplementaryFig.8)inirradiatedGNSsamples,pos-
siblyduetotheevolutionfromfaultedloops(e.g.,1/3<111>loopsin
austenite)tounfaultedloops(e.g.,1/2<111>loopsinmartensite).Itis
interesting to notice this trend is very different from that both the
density and size of radiation defects significantly increase with
increasingdosewithintheearlystageinirradiatedCGcounterpartsin
this work, as well as in other irradiated materials (including nanos-
tructuredandCG)reportedinpreviousworks25–30.
Since the formation of radiation defects depends on material
compositions and radiation conditions, and an enhanced radiation
resistanceistypicallyreflectedbyalowerdefectdensity,thus,theratio
ofradiationdefectsdensityinaNCmaterialtothatinitsCGcoun-
terpart(i.e.,specificradiationdefectsdensity,ρ /ρ )underasame
NC CG
irradiationprocesscanbeusedtoreflecttheincrementofradiation
resistancebyformingnanostructuresindifferentalloys.Asshownin
Fig.2dandSupplementaryFig.9,thespecificradiationdefectsdensity
inGNS304NGismuchlowerthanthoseinotherkindsofNCmaterials
underasimilarradiationcondition.Althoughtheratioinnano-twinned
Agisalsolow(~0.22)atalowdose(0.12dpa),itincreasesrapidlywith
increasing dose26, possibly due to its low microstructure stability
underirradiation.Incomparison,thespecificradiationdefectsdensity
ofGNS304NGdecreasesgraduallywithincreasingdoseandreaches
~0.038at155dpa,evenlowerthanthatinT91steel,whichiswidely
used in structural components (suchaspressurevessels,pipes and
powerboilers)incurrentnuclearreactors27.
Inadditiontoasignificantlyenhancedradiationresistancetothe
formation ofDLsatroomtemperature,anenhancedradiationresis-
tance to the formation of voids, which are also a typical kind of
recombined radiation-induced defects, has been observed in GNS
304NG samples at high temperatures. As shown in Supplementary
Fig.10,alargenumberofvoidsformsinCGsamplesafteraradiation
doseof~80dpaat450°C,whilevoidscanberarelyobservedinGNS
counterparts.

So, what mechanisms enable GNS 304NG steel to achieve such an
excellent irradiation tolerance? One should be its specific nanos-
tructureswithahighstability.Conventionally,refiningmicrostructure
leadstoanenhancedradiationresistance,mostlyduetotheexistence
ofnumerousGBsactingassinksofradiation-induceddefects10.How-
ever, the instability of refined microstructure is an Achilles heel of
traditional NC materials under irradiation, mostly due to the high
storedenergyofGBsinthem.Inthiswork,Σ3(111)GBsaremassively
formedintheGNSsurfacelayer.Theyaretypicallyverystableunder
heating or irradiation, mostly due to the low stored interfacial
energy10,31. Meanwhile, numerous SFs and partial dislocations are
formedandpossiblyleadintoadistinctrelaxationofGBs32.Andthey
might also promote the annihilation of radiation-induced defects
insidegrainsandreduceGBmigrationunderirradiation(asdiscussed
later).Moreover,thedislocationlinesforma3Dnetwork,whichmight
reduceGBmigrationnotonlybyabsorbingradiation-induceddefects
insidegrainsbutalsobyinteractingwithGBsandformingstrangle-
holdonthem33.Therefore,suchaspecialmicrostructure(combining
low-energy interfaces, SFs, and DNs) contributes to a high stability
fromboththermodynamicandkineticperspectives.Andthispointhas
alsobeenconfirmedbyMDsimulations(seeSupplementaryNote1).
Therefore,thegrainsizeisstillwithinnanometerscaleataveryhigh
radiationdose(e.g.,86±26nmafter155dpa,seeFig.2bandSupple-
mentary Fig. 6) at room temperature, or even after irradiation at a
high-temperature(e.g.,72±32nmafter80dpaat450°C,seeSupple-
mentary Fig. 10). Meanwhile, the GBs are still mainly composed of
low energy onessuch as Σ3 interfaces and LAGBsinirradiatedGNS
samples(SupplementaryFig.11).
Asmentionedearlier,theDLsdensitysignificantlydecreaseswith
increasingdosefrom2to10dpainthepeakdamageregionsindif-
ferentirradiatedGNSsamples(seeFig.2candSupplementaryFig.8).
Inaddition,thedecreasingDLsdensitywithincreasingdosewasalso
observedinasameirradiatedGNS304NGsample(seeSupplementary
Fig.12).Incontrast,theDLsdensityincreaseswithincreasingdosein
CG counterparts in this work, as well as in other materials under
irradiation25–30. Therefore, we believe that there should be a more
importantanti-radiationmechanisminGNS304NG.
MartensitictransformationoccursinbothGNSandCGsamples
under irradiation, but with totally different features. In the peak
damage region in GNS samples, the fraction of martensitic phase
increases from ~10% to ~50% after an irradiation of 5dpa (Supple-
mentary Fig. 13), and reaches almost a complete distribution after
12dpa(SupplementaryFig.6)atroomtemperature.Incontrast,alarge
fractionofausteniticphaseisstillretainedinthepeakdamageregion
in irradiated CG samples even after a much higher dose of 155dpa
(Fig.2a).Therefore,themartensitictransformationoccursmuchfaster
andinabroaderregioninGNSsamplesthaninCGsamplesintheearly
irradiationstage.Furthermore,X-raydiffraction(XRD)analysesreveal
thatthemartensitecontentincreasesgraduallywithdoseinirradiated
GNS samples, while it is almost constant in irradiated CG samples
(Fig.3aandSupplementaryFig.14).ItisnotedthattheXRDresults
shouldreflectthemeanphaseinformationinasurfacelayerof~5μmin
thickness(largerthantheirradiatedthickness)duetotheabsorption
ofX-rayinthesteels34,sothatthemartensitevolumefractionvalues
obtainedbyXRD(seeFig.3a)arelowerthanthoseobtainedbySAED
and TKD in peak damage regions (see Fig. 2a and Supplementary
Figs. 6 and 13).Combining the results obtained by SAED, TKD, and
XRD, we can see that the martensitic phase transformation should
occuronlyinthepeakdamageregioninirradiatedCGsamplesatroom
temperature(SupplementaryFig.6i),whileitshouldoccuralmostin
the entire irradiated region in GNS counterparts (Supplementary
Fig.6j).Moreover,almostnomartensitetransformationisobservedin
irradiatedCGsamplesathightemperatures(e.g.,450°C),whileitstill
occurs vastly in irradiated GNS counterparts, as demonstrated by
comparingthecorrespondingSAEDpatternsinSupplementaryFig.10.
Therefore, an unexpected novel large-scale adaptive martensitic
transformation (AMT) from the peak damage region to the sur-
roundingregionsinGNS304NGunderirradiationcanbesuggestedfor
the 1st time. In brief, the AMT process is different from the conven-
tional martensitic transformation process in CG austenitic stainless
steels under irradiationin thatthe AMTprocess occurs faster, pro-
gressesadaptivelyfromthepeakdamageregiontothesurrounding
regions,andcanbeactivatedathightemperatures.
Mostlyduetotheirlowerpackingdensities,bodycenteredcubic
(BCC)metalsaretypicallymoreresistanttoradiationdamagesthan
facecenteredcubic(FCC)metals1,10,35.Therefore,theradiationresis-
tance is expected to be enhanced in GNS 304NG after martensitic
transformation. In this work, an unexpected contribution is further
provided to the extraordinary radiation resistanceduring the large-
scale AMT process itself. As revealed in Supplementary Fig. 15, in
comparisonwithGNS304samples,whichpossessasimilargrainsize
distribution but much larger initial martensite fraction36, the DLs
densitiesinGNS304NGsamplesareloweratdoses>6dpa,although
they are higher at ~1dpa mostly due to a much smaller martensite
fraction.
In fact, radiation induced martensitic transformation has been
observed insuchmaterialsasausteniticstainlesssteels18,19,titanium
alloys37,andTRIPalloys20previously,butmuchhigherradiationdoses
arerequiredtoinitiatethetransformation.Inthiswork,asignificant
increaseinmartensitecontentisobservedinGNSsamplesunderlow
radiation doses of 1−2dpa at room temperature. Further character-
izations by high-resolution STEM show that the latticearrangement
changes near SFs at first (Fig. 3b, c), implying that the interactions
between pre-existing SFs and radiation-induced defects boost the
martensitictransformation.Therefore,theexcessfreeenergiesrelat-
ingwithSFsandradiation-induceddefectsmightprovideaprimary
driving force for martensitic transformation, and the existence of
numerousSFsshouldpromotetheAMTprocesswidelyinGNS304NG.
Meanwhile,nochemicalsegregationorvariationcanbedetectedafter
phase transformation (Supplementary Fig. 16), confirming that the
radiation inducedphasetransformationshouldbeaprocesswithout
long-rangemassdiffusionorimplantedions(duetoradiation).Further
MD simulations (see Supplementary Note 1) demonstrate that the
mechanismofmartensitictransformationisdominatedbyinteractions
betweenSFsandradiation-induceddefectsinirradiatedGNS304NG.
And transformation is a 2-step process, conforming to the classical
Kurdjumov-Sachsrelationship.ItisnotedthatonlytheBCC(α’)mar-
tensitic phase was observed by both TEM and XRD in this work,
althoughthehexagonalclose-packed(HCP,ε)martensiticphasemight
also be evolved from overlapping bundles of SFs on {111} planes in
principle38,39.Sincelatticerearrangementoccursduringphasetrans-
formation and mightconsume a large number of radiation-induced
defects, the radiation resistance is thereforesignificantly enhanced.
ThisagreeswiththattheGNS304NGsamplesshowhigherradiation
resistancethanGNS304counterpartsatdoses>6dpa(Supplementary
Fig.15),whentheAMTmechanismisactiveintheirradiationaffected
rangeinGNS304NG(seeFig.3a),butthefractionofformedmarten-
siteisstillmuchlowerthanthatinGNS304(near100%).
Typically,highercompressiveresidualstressesareintroducedin
thesurfacelayerbySMRT36,whichmightpromotetheAMTprocessby
decreasing the energy barrier of martensitic transformation. In this
work, the large-scale AMT and enhanced irradiation resistance still
occurinGNSsamplesathighertemperatures(e.g.,450°C),whenthe
residualstressesarealmostcompletelyreleased.Therefore,itshould
beSFsplayingthedominantroleintheAMTprocessinGNS304NG.
ThispointhasbeenclarifiedbyMDsimulationsofAMTmechanism
(Supplementary Note 1). However, it is noted that stresses were
introduced to accelerate the martensitic transformation during MD
simulationsinthiswork.Andchangesinresidualstressesmightoccur
after irradiation at room temperature. Therefore, further investiga-
tionsareexpectedtoclarifytherelationshipbetweenresidualstress
and the AMT process under irradiation, especially at room
temperature.

Generally, the new phases are isolated and in-homogeneously dis-
tributedintheparentphasesinnormalausteniticstainlesssteelsafter
radiationinducedmartensitictransformation.Theywillcausestress
concentration and result in significant degradations in mechanical
properties, e.g., ductility, fatigue, and creep resistance18,19,37,40. As
shown in Fig. 4a–c, irradiated CG samples exhibit typical radiation
inducedsurfacehardening,i.e.,thehardnessismarkedlyraisedatthe
topsurface(from~2.31GPato~4.37GPa,ΔH/H ~0.89,accordingtothe
I
nanoindentationresults)andthendecreasessharplyalongdepthafter
radiation.Incomparison,althoughthesurfacehardnessofGNSsam-
ples is also raised (from ~4.93GPa to ~6.45GPa, ΔH/H ~0.31) after
I
radiation,thehardeningtrendislowerandoccursinamuchbroader
region.Thisshouldbebenefitedbytheexcellentcoordinateddefor-
mationabilityofGNS36,41,42,aswellasthelarge-scaleAMTinthesurface
layer.Frompreviousstudies,thesharpsurfacehardeningistypically
prone toinducesurface embrittlement43,44,whilethe gradienthard-
ness distribution in a thicker surface layer is expected to result in
superiormechanicalproperties41,42,45.
As demonstrated by micropillar compression tests (Fig. 4d, f),
distinct slip traces form on GNS samples after irradiation at room
temperature,whicharesimilartothoseformonGNSsamplesbefore
irradiation,andnofractureoccursafteracompressionstrainto~22%.
TheseindicatethatnosurfaceembrittlementoccursinirradiatedGNS
samples.Sinceahighfractionofmartensitegrainshasformedinthe
surfacelayerofirradiatedGNSsamples,wecaninferthattheradiation
inducedlarge-scaleAMTprocessshouldnotdegradethemechanical
properties of 304NG. In fact, outstanding mechanical properties,
including superior strength-ductility synergy and fatigue resistance,
havebeenobservedinGNSausteniticstainlesssteels(316Land304)
previously, in which the surface layer is mostly composed of
deformation-induced martensitic nano-grains36,46. Thus, this work
indicatesthespecialmicrostructureofGNS(includingnumerousΣ3
GBs,SFs,andDNs)exploresabeneficialmartensiticphasetransfor-
mation under irradiation, which not only decreases the number of
radiationdefectsbutalsoimprovesthemechanicalpropertiesofsteels
servedinreactors.
In summary, a strategy to develop radiation resistant steels
with high design flexibility could greatly promote the safe appli-
cation of nuclear energy. In this work, we discovered a novel
radiationinducedlarge-scaleAMTmechanisminGNS304NGwith
a high density of SFs, Σ3 GBs, and LAGBs (prepared by surface
nano-crystallization technologies47 which are low cost and pos-
sesshigh designflexibility,seeSupplementaryFig.18).TheAMT
mechanism keeps effective under irradiation from low doses to
extremely high doses, from low temperatures to high tempera-
tures. The introduction of large-scale AMT changes the detri-
mentalmartensitictransformationintoabeneficialprocessunder
irradiation for the first time, which develops a steel with strong
radiation resistance. Therefore, the surface nano-crystallization
technique and AMT mechanism present brilliant application
potentialities in reactors, especially for core components under
high irradiation doses and dose rates. Furthermore, since radia-
tioninducedmartensitictransformationmightoccurindifferent
metals such as austenitic stainless steels, titanium alloys, and
TRIP alloys, this work provides a novel way to develop super-
metals with high radiation resistance to be used in current and
futurenuclearfacilities.

Methods
Inthiswork,anucleargradeausteniticstainlesssteel304(304NG)was
studied.As-receivedplatesampleswereinacold-rolledstatewitha
microhardness of ~2.0GPa. Its chemical composition was (in wt.%)
19.51Cr,8.87Ni,1.73Mn,0.27Si,0.021P,0.04C,andbalanceFe.The
initialmicrostructureconsistedofequiaxedgrains(around93%aus-
teniteand7%martensite)withanaveragesizeof~26±10μmandfew
deformationtwins(SupplementaryFig.1).Inaddition,fewdislocations
andSFsexistedinthegrains.
GNSlayerwithathicknessof~700μmwasin-situproducedonthe
surface of 304NG plates by surface mechanical rolling treatment
(SMRT).TheprocessingofSMRThasbeenillustratedinSupplemen-
taryFig.2a.DuringSMRT,apolishedWC/Coballof8.0mmindiameter
waspressedintotheplatesurfacewithadepthofa androlledonthe
p
surfacealongtheydirectionatavelocityofv.Subsequently,theball
y
movedalongthexdirectionataperstepdistanceofP andrepeated
x
therollingprocess.TheprocessingparametersofSMRTinthepresent
workarelistedinSupplementaryTable1.ItisnotedthatSMRTisa
methodimplementingsevereplasticdeformationatthematerialsur-
face.However,SMRTtypicallyapplieshigherstrainandstrainratesin
thesurfacelayer,anditshowsadvantagesoflowcostandhighflex-
ibility in processing large-sized components, compared to conven-
tional severe plastic deformation methods such as equal channel
angularpressingandhigh-pressuretorsion1,2,9.
After SMRT, these samples were cut into squares
(10×10×1.9mm3) by wire electrical discharge machining (WEDM).
Andthenelectrolyticpolishing(in10%perchloricacidalcoholsolu-
tion,20V,−20°C,45s)wasusedtoremoveasurfacelayerof~2μm
inthicknessfromthetreatedsurface,toexcludeanypossiblesurface
contamination during the sample processing. Finally, GNS 304NG
sampleswereobtained,asrevealedinSupplementaryFig.3.Mean-
while, as-received samples were also cut into squares
(10×10×1.9mm3)byWEDM.Subsequently,theyweresuccessively
polishedbyusing400,800,and2000gritemerypapers,followedby
electrolyticpolishing(in10%perchloricacidalcoholsolution,20V,
−20°C,90s).Thus,CGcounterpartswereprepared.

TheFeionradiationofGNS304NGsamplesandCGcounterpartswas
carriedoutattheionimplantationterminalintheLowEnergyHeavy
Ion Accelerator Facility (LEAF) platform, at the Institute of Modern
Physics,ChineseAcademyofSciences,asschematicallyillustratedin
Supplementary Fig. 2b. The LEAF could perform homogeneous ion
implantation into the sample surface layer within an area up to
50×50mm2 by Fe ion implantation. Thus, up to approximately 25
squared samples were irradiated simultaneously under the same
experimental conditions. Firstly, Fe13+ ions with an incident particle
energyof0.69MeVu−1wereimplantedintothesamplesatafluxof
2.5×1011ionscm−2s−1.Pre-experimentswereperformedwith3differ-
ent ion fluences, i.e., 5.0×1015, 1.0×1016, and 5.0×1016 ionscm−2,
respectively.Theradiationinducedlarge-scaleAMTwasdiscoveredin
all these pre-experimental samples. Then, in order to confirm and
elucidate this mechanism, we conducted a series of validation
experimentswithdifferentionfluxes,energies,andtemperatures.For
example, Fe10+ ions with an incident particle energy of 0.62MeVu−1
wereappliedatafluxof2.8×1011ionscm−2s−1,andthesampletem-
perature was controlled to ~80°C through water cooling. The total
implanted Fe10+ ion fluences were within 6.50×1014 to 1.625×1017
ionscm−2.
High-temperatureradiationexperimentswerealsoconductedat
350,400,and450°C.Forexample,duringradiationat450°C,Fe15+
ionswithanincidentparticleenergyof0.65MeVu−1wereappliedata
fluxof7.0×1011ionscm−2s−1.Themaximumionfluencewas8.0×1016
ionscm−2.

Thedistributions ofradiationdose and implanted Feconcentration
(C )alongdepthfromtheirradiatedsurface(w)werederivedfrom
Fe
DoseðwÞ=φ(cid:2)vðwÞ=M ð1Þ
and
C ðwÞ=D(cid:2)RðwÞ ð2Þ
Fe
whereφwastheincidentionfluence,v(w)wastheaveragevacancy
generatedperunitdistanceatwforeachincidention,R(w)wasthe
average atomic implantation ion concentration distance at w, M
was the atomic number density of the material, D was the unit
thickness.Here,v(w)andR(w)wereobtainedfromtheSRIMout-
putfiles(Vacancy.txtandRange.txt,respectively)byusingaSRIM-
2013softwarepackage,whichwasusedtoevaluatehowionslose
energy into matter and the final distribution of these ions after
they stop withinthetargetbasedontheMonteCarlomethod48.In
the present work, we used the Kinchin-Pease model in the SRIM
simulationtoevaluateFeionirradiation(34.68MeV)inamaterial
simulating 304NG (72at.% Fe, 19at.% Cr, and 9at.% Ni), with a
nominal displacement threshold energy of 40eV for each
element49.
ThedistributionsofradiationdoseandFeconcentrationalong
depth obtained in the validation experiments are given in Supple-
mentary Fig. 4. And the results of the pre-experiments and high-
temperatureexperimentsarepresentedinSupplementaryTable2.

Scanningelectronmicroscopy(SEM)observationswereconductedon
anFEIVerios460unitundertheconcentricbackscatteringmode.TEM
andSTEMcharacterizationswerecarriedoutinanFEITalosF200X
microscopeoperatedat200kV.Andhigh-resolutionSTEMobserva-
tionswerecarriedoutonaTitanThemisG2-300doubleaberration-
correctedunitoperatedat300kV.
Toacquirethephaseandinterfaceinformationinpeakdamage
regions in GNS 304NG samples before and after irradiation, TKD
analysiswascarriedoutonaZEISSGemini300SEMunitoperatedat
30kVwithacurrentof6.0nA.Inaddition,electron back-scattering
diffraction(EBSD)analysiswasperformedonas-receivedCGsamples
usinganFEINova450SEMunitoperatedat18kV.Thescanningstep
sizewassetwithin1/5to1/10oftheaveragegrainsizeinEBSDandTKD.
The EBSD and TKD data was analyzed by HKL Channel 5 (Version:
5.12.67.0).
Cross-sectional SEM and EBSD samples were firstly electro-
depositedwithaprotectiveNicoating,subsequentlycutbyWEDM,
mechanicallypolished,andfinallyelectrochemicallypolishedin10%
perchloricacidalcoholsolution(20V,−20°C,90s).
Cross-sectional TEM foils were prepared by focused ion beam
(FIB) using an FEI Nanolab Helios 650 (Ga+) unit. A Pt layer was
deposited to protect the sample surfacebefore FIBmilling. Asame
direction(alongtheSMRTprocessingdirectionx)wassectionedforall
samplestoexcludethepossibleeffectsoftexture.AndTEMsamples
werefinallypolishedatlowvoltage(2kV)andlowcurrent(39pA)to
cleanthepossibleFIB-induceddamages.

Inthiswork,themainradiation-induceddefectswereDLsinsamples
irradiatedatroomtemperature.InirradiatedCGsamples,although
therewerealreadyfewdislocationsandSFsbeforeirradiation,and
dispersive martensite grains were formed after irradiation, DLs
couldstillbeclearlydistinguishedbyWBDFTEMimaging.Inirra-
diatedGNSsamples,thepre-existenceofnumerousotherkindsof
defects(e.g.,dislocationtanglesandstackingfaulttetrahedrons,
typicallywithsizes<3nm)mightbringmistakesinlabelingradiation
inducedDLs.Sincepreliminaryanalysesindicatedthatthesizeof
DLs in GNS 304NG was mostly within the range of 3–6nm (see
Supplementary Fig. 7), we used the following criteria to distin-
guishDLs:
1) BoththeBFandtheWBDFTEMimageswereanalyzed.
2) Mark the 1/2a<111> DLs according to the directions of related
Burgersvectors.
Subsequently,theDLsdensityanddiameterwereobtained.Here,
theDLsdensitywascalculatedfromthenumberofDLs(N)andthe
volume of observed zone (V) in a TEM foil, i.e., N/V, where V was
calculated from length×width×thickness (within 70–110nm, mea-
suredinFIB).Theresultswereaveragedfrom3differentTEMfoilsfor
each irradiated sample. 1–2 images were taken from each foil. It is
notedthatalmostnoDLsshouldbeintroducedintoTEMfoilsduring
FIBpreparation,asrevealedbycarefullycheckingthenon-irradiated
sampleinSupplementaryFig.5.

While XRD analysis typically provides phase information in sam-
ples with a larger dimension than TEM does, irradiated samples
were analyzed on a Bruker D8 Advance XRD diffractometer
(16kW, using Cu Kα radiation) with a scanning angle range of
40–100° and a scanning rate of 2θ=2°min−1. The fraction of a
phase (i.e., γ-austenite or α‘-martensite) was determined from
XRDpattern50 accordingto:
1 P n I i j
VðiÞ= P n j=1 P R i j ð3Þ
1 n Iγ j + 1 n Iα j 0
n j=1Rγ j n j=1Rα j 0
wherenwasthenumberofpeaksofthephaseusedincalculation,Ij
i
andRjweretheintegratedintensityandmaterialscatteringfactorfor
i
thejpeakofthephasei,respectively.

The distribution of microhardness along depth was measured on
cross-sectional samples, by measuring microhardness at different
depths from the irradiated surface. A Qness Q10A+ microhardness
tester equipped with a Vickersindenterwas used,with amaximum
loadof20×gandadwellingtimeof10s.Eachmicrohardnessvalue
wasaveragedfromatleast5measurements.
Meanwhile, nanoindentation tests were conducted on sample
surfacefromaplanerviewbyusinganAgilentG200unit(Berkvoich
indenter, strain rate 0.05s−1). The maximum indentation depth was
2000nmduringatest.Theloaddisplacementdataobtainedduring
loadingwereanalyzedusingacontinuousstiffnessmodetoobtainthe
indentation displacement dependence of hardness51. Subsequently,
theequivalentsurfacehardnessofasamplewasestimatedfromthe
indentationdisplacementdependenceofhardnessaccordingtothe
Nix-Gaocorrectionmodel52.Atleast5measurementswereperformed
onasample.
Moreover,in-situmicropillarcompressiontestswereperformed
byaFemto-Toolsin-situnanoindentationequippedinaZeissGemini
SEM460unit.Micropillarsamples(witha~3°taperangle)of~5μmin
theaveragediameterand~10μmintheheight(i.e.,withaheight-to-
diameterratioof2)werepreparedalongtheirradiationdirectionfrom
sample surface by FIB. The compression process was loaded at a
constantdisplacementrate(5nms−1)upto~22%ofengineeringstrain
undera10μmSi-basedflat-punchindenter.Theforce-displacement
datawascollectedduringcompressiontesting,andareal-timevideo
was applied to record the deformation process of the compressed
micropillar.

An atomic model of Fe-Cr-Ni alloy (71.9at.% Fe, 18.5at.% Cr, and
9.6at.%Ni)withgradientnanostructurewassetupasGNS304NGideal
model for MD simulations. Based on our experimental results, the
atomic model was set up with a box volume of
34.73×30.08×40.00nm3, which contains gradient structure con-
sistedof4consecutivelayers(inanFCCarray)withincreasingthick-
nessfromtoptobottom,similartoexperimentalsamples.Between
neighboringlayers,Σ3(111)GBswerebuiltbymakingthez-directionof
everygrainat[111]orientationfollowingcoincidencesitelatticethe-
ory. In every layer, there were low-angle GBs between neighboring
grains.
Tofurtheranalyzethemechanismoflarge-scaleAMT,webuiltthe
2nd MD model with a box volume of 8.96×9.03×9.58nm3. In this
model,tworegionswereset,includingasphericalBCCphaselocated
at the center of model, surrounded by an FCC Fe-Cr-Ni alloy. Fur-
thermore,toexploretheroleofSFsonextensionoflarge-scaleAMT
fromcentralBCCphasetosurroundedFCCphase,twocrossSFswere
pre-builtinFCCphasebutendedatBCC/FCCinterface.

Bychoosingthisatomicmodel,werelaxedthesamplebyconjugate
gradientalgorithm,whichwasfollowedbythermalrelaxationtothe
targettemperature(300K)togetamoreenergeticstablestate.
Afterconjugategradientrelaxationandsubsequentheating,fol-
lowingtheclassicalcascadesimulationmethod7,aprimaryknock-on
atom(PKA)wasselectedrandomlyfromthetopmostlayersandwas
given a kinetic energy to induce further cascade in the system to
explore the irradiation tolerance of GNS 304NG model a (cid:1) lloy. The
energy of PKA was 80keV and emitting direction was [001], which
couldwellsimulatetheionimplantationexperimentsafterconsider-
ing the Coulomb screening effects as done in this work. A
Nosé–Hoover53thermostatwasappliedtotheoutermostlayerofthe
box with a thickness of approximately 3 atomic layers to regulate
atomicvelocitiesbasedonthetargettemperature.Thus,duringcas-
cade simulation, the system was simulated with a microcanonical
ensemble(constantparticlenumber,constantvolume,andconstant
energy)andcanonicalensemble(constantparticlenumber,constant
volume, and constant temperature) to the interior and boundary
atomsofthesimulationbox,respectively(SupplementaryFig.20).
ConsideringtheeffectsinducedbyresidualstressafterSMRTand
stresses induced by phase transformation and ion implantation, we
firstly carried out tensile simulations with 0.1Åps−1 strain rate to
acceleratethemicrostructureevolutionaftercascadesimulations.For
thetensiledeformationsimulations,theappliedforcewasalongthez
direction. Furthermore, we also carried out shear simulations with
0.1Åps−1strainrate,whiletheappliedforcewasalongthexdirection.
Tworegionswiththicknessof20Åfrombothsurfacesofthesimulation
boxesweresetasrigidbodies,resultinginafixedboxshapeunder
theinfluenceofthetensileorshearforce.Theotherregionsofthe
boxesweremobileduringthewholesimulationprocess.Forthe2nd
MDmodelasdescribedabove,thetensilesimulationswith0.1Åps−1
strainratealongxdirectionwerealsoperformedafterfullConjugate
Gradientrelaxations.
Inboththeheatingandloadingprocesses,atomicpositionsand
velocities were integrated and updated at each time-step (1fs) to
maintainaconstanttemperature.Theinteratomicpotentialwasbased
on the semi-empirical model developed with the embedded atom
method,areliableclassicalatomisticmodelforsimulationsofcrystal
defectsintheFe-Cr-Nisystem54.Withpolyhedraltemplatematching
method embedded in an atomic analysis and visualization software
calledOvito(version:3.11.3)55–58,thephasesdistribution,SFs,andGBs
could be distinguished in the simulation model. Here, SFs were
recognized as HCP arrays due to their similar lattice arrangements,
which was distinguished by further surrounding lattice analysis. All
simulations were carried out using the MD package Large-scale
Atomic/MolecularMassivelyParallelSimulator(version:2Aug2023)59.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "This work provides interesting results concerning radiation tolerance improvements via surface treatment. The ion irradiation is an acceptable starting point for further investigation of radiation effects in this unique microstructure. Generally the conclusions are supported by the data, and the experimental and computation methods are sound. However, additional detail and explanation of certain concepts is suggested. This work is recommended for minor revision, then publication pending that revision.", "reviewer_2": "The draft “Extraordinary Radiation Resistance in Metals via Adaptive Martensitic Transformation” by authors presents the findings of an experimental investigation demonstrating the formation of unexpectedly lower densities of dislocation loops and voids in a designed nanocrystalline austenitic stainless steel under ion irradiation. Accordingly, they find that a novel and interesting mechanism called adaptive martensitic transformation (AMT) contributes to the superior radiation resistance. And they further clarify and confirm the findings using MD simulations. Meanwhile, they also find superior mechanical properties in the irradiated nanocrystalline samples.\nThe experimental design is reasonable. The microscopy studies are thorough and the mechanical measurements are credible. And the process to produce the designed nanocrystalline structure for AMT in mm-scale thicknesses at low cost is desirable. In the referee’s opinion, this draft can be accepted for publication after several issues be addressed.\n1. Martensitic transformation is a topic worthy of in-depth research, such as TRIP effect, shape memory alloys, magnetism, etc. Residual stress has a significant impact on the martensitic transformation. Although the authors mentioned the release of residual stress at high temperatures in the line 261-267, the effect of residual stress on the AMT mechanism at room temperature deserves further investigation and may become a means of surface modification. Do the authors pay attention to the relationship between martensitic transformation and residual stress in the irradiation samples?\n2. The manuscript claims that the AMT mechanism is related to high-density SFs. So, what is the value of it. A detailed value would be appreciated.\n3. The authors focus on studying the microstructure of the peak damage region. According to the dose and ion distribution map provided in Supplementary Fig. 3, the irradiation dose and ion concentration in the peak damage area are much higher than other irradiation regions. So, will this affect the martensitic transformation in this work?\n4. In lines 105-106, the authors mention a 3D dislocation network was formed in 304NG by SMRT. That’s interesting. However, there seems to be no experimental evidence, and only been demonstrated through MD simulations.\n5. In lines 141-143, “it is interesting to notice that both the density and size of DLs significantly decrease……0.5 to 5 dpa……” the range of the dose is incorrect (refer to Supplementary Fig. 4).", "reviewer_3": "This manuscript states the radiation tolerance of metals can be enhanced by a mechanism called adaptive martensitic transformation. The manuscript uses the experimental methods and molecular dynamics simulations to reveal this mechanism. After reviewing this manuscript, the reviewer has some comments. See details in below.\n1. Is the first paragraph the abstract? If this is an abstract, we should not cite the references. Please remove these references and reorganize the references in the whole manuscript.\n2.This manuscript is comparing the radiation behaviors of the CG and GNS material, I would suggest revealing the grain size distribution for both CG and GNS. The manuscript only listed the average grain size of CG is 26 µm, what is the grain size and its distribution for GNS?\n3.This manuscript discussed the phase transformation under irradiations. What about the radiation induced grain growth, especially under the high dose after 155 dpa? The authors should also reveal the grain size after 155 dpa and explain why the material is still within nanometer scale after high irradiation dose in more details. I can see the authors mentioned combining low energy interfaces, SFs, and DNs contributes to a high stability from both thermodynamic and kinetic perspectives. Please explain more how these interfaces, SFs and DNs will contribute the structure stability in thermodynamic and kinetic perspectives.\n4. I suggest all the SAED patterns in this manuscript should be identified. Especially the diffraction rings for the martensite in the GNS images.\n5.In Figure 2, this manuscript is discussing the DL size of the CG and GNS sample. However, the error bar for the CG sample is way bigger than the GNS sample. Can you explain this?\n6.Based on the discussion and Figure 2, it looks like after 10 dpa, the DL size and density remain flat. Please explain why 10 dpa is the critical irradiation point in more details? Why will the DL size not increase, and DL density will not decrease after 10 dpa? Is this related to the martensitic phase?\n7. Please check all the references and keep them in the same format. Especially the format for the journal’s name.\n8. In extended data figure 3, why did not put the images for 155 dpa? It only includes the images up to 85 dpa. Since 155 dpa is the final dose, I would suggest adding the images for 155 dpa.\n9. In extended data figure 4, the figure 4 (b) has no scale bar. Also, it looks larger of the DL in Figure d than Figure c. Please clarify this.\n10. In extended data figure 6, it says the white arrows show g= [001]. However, it is very difficult to see in the images, please improve the image quality.\n11. In the supplementary Figure 3, I would suggest changing the curve colors in the right image to make it more obvious to see the difference."}, "author_response": "Response-to-Reviewers:\nManuscript ID: NCOMMS-25-38174-T\nTitle: Extraordinary Radiation Resistance in Metals via Adaptive Martensitic\nTransformation\nAuthors: S. Zhang et al.\nThe authors would like to thank all the reviewers for the valuable comments and\nsuggestions which are definitely helpful for making the manuscript more accurate and\nclearer. According to the comments and suggestions, we have performed some\nadditional works (including some new experiments, e.g. see the replies to Reviewer 2:\nComment 1; Reviewer 3: Comment 9), and expanded the introduction and discussion\n(including the TRIP effect, see the reply to Reviewer 2: Comment 1). And we have\nrevised the manuscript accordingly, as highlighted in yellow in the revised manuscript.\nBelow detailed responses to the comments are listed.\nReviewer #1:\nThis work provides interesting results concerning radiation tolerance improvements via\nsurface treatment. The ion irradiation is an acceptable starting point for further\ninvestigation of radiation effects in this unique microstructure. Generally the\nconclusions are supported by the data, and the experimental and computation methods\nare sound. However, additional detail and explanation of certain concepts is suggested.\nThis work is recommended for minor revision, then publication pending that revision.\nREPLY：Thanks for your insightful comments and constructive suggestions!\n1) P2 Line 67: What is the different between gradient nanostructured material and\nsevere plastic deformation at the surface of the material?\nREPLY: In this work, we used a surface severe plastic deformation method, surface\nmechanical rolling treatment (SMRT), to produce gradient nanostructured 304NG.\nFrom this aspect, SMRT is a method implementing severe plastic deformation at the\nsurface of materials. However, higher strain and strain rates are typically required at the\n at the\nmaterial surface in SMRT than that in conventional severe plastic deformation methods\n[1, 2] to achieve gradient nanostructures, as revealed in Ref. [3]. Moreover, SMRT\nshows advantages of low cost and high flexibility in processing large-sized components\ncompared to conventional severe plastic deformation methods such as equal channel\nangular pressing [1] and high-pressure torsion [2]. Meanwhile, in addition to surface\nsevere plastic deformation methods, some other methods, such as electro-deposition [4]\nand magnetron sputtering [5], can also be applied to produce GNS materials.\nFollowing your question, the difference between gradient nanostructured material and\nsevere plastic deformation at material surface has been mentioned in the revised\nmanuscript (P14 Lines 390-395).\n2. P3 Line 74: What is the reason that 304 specifically was chosen for this study? What\nwould the best application of this technique in reactor be? Is it more advantageous to\na component which received a high dose rate, or one that receives dose, but is more\naffected by temperature?\nREPLY: Mostly due to its outstanding performance of mechanical properties, high-\ntemperature resistance, corrosion resistance, and acceptable radiation resistance, 304\nstainless steel has been applied for decades as the primary components in pressured\nwater reactors (PWR) and boiling water reactors (BWR) [6]. For example, this material\nhas been used as the core structures, primary pipes, pump materials and so on in PWRs,\nand used as plates of core shroud, core supporting structures, pipes in jet pumps and so\non in BWRs. Furthermore, 304 stainless steel has also been considered as a main\nstructural material in both nuclear fusion reactors [7] and next-generation reactors [8].\nTo understand and improve their radiation resistance would thus be useful and\nnecessary for the current and future nuclear reactors. Therefore, we selected 304\nstainless steel in this work, to further improve its resistance to radiation damage.\nSMRT and the AMT mechanism can be applied on pipes and plates of 304 stainless\nsteel in reactors as mentioned above.\nAccording to the present results, the AMT mechanism in GNS-304 NG can be applied\nafter high dose radiation, e.g. up to 150 dpa. In fact, the radiation dose rate in ion-\nimplantation experiments was also high (up to 2.7810-4 dpa s-1) in this work. Thus,\nthis mechanism can also work for cases under a high radiation dose rate. As for the\ntemperature effect, we have proved that this mechanism still works well at temperature\nup to 450 C, while beyond which no experiments have been performed at present.\nTherefore, the AMT mechanism should have advantageous to a component which\nreceives a high dose and/or a high dose rate at temperatures at least up to 450 C.\nIn the revised manuscript, the reason that 304 stainless steel was chosen has been added\nin P3 Lines 80-82, and the application possibilities of SMRT and AMT mechanism have\nbeen further discussed in P13 Line 372-374.\n3. P6: Line 144: The difference in loop size and density evolution could be understood\nin the context of faulted loop growth and unfaulting, which could lead to decreasing\nloop size\nREPLY: Yes. The evolution from faulted loops (e.g. 1/3<111> loops in austenite)\nbefore martensitic transformation to unfaulted loops (e.g. 1/2<111> loops in martensite)\nafter martensitic transformation can be used to understand the difference in loop size\nand density evolution in this work, especially for the decrease of loop size shown in\nFig. 2c.\nFrom Fig. 2, Fig. 3 and Supplementary Fig. 7 (in the revised manuscript, hereafter the\nsame), one can see that before martensitic transformation dominates the defect\nevolution process (e.g. at doses below 2 dpa), the faulted 1/3<111> loops are primary\ninterstitial cluster defects in FCC austenite, which have low diffusion capability [9] and\ntheir average size could increase through absorption of interstitials induced by\ndisplacement cascades. When the martensitic transformation dominates the defect\nevolution process (e.g. at doses larger than around 2 dpa), the formed loops in the\ntransformed martensitic phase (BCC) are unfaulted 1/2<111> loops with very low\ndiffusion energy barrier [9], resulting in their fast absorption by a high density of grain\nboundaries, dislocation networks and stacking faults, which would also absorb the\nradiation induced point defects. Meanwhile, the AMT mechanism further annihilates\nthe point defects, resulting in a decreasing number of interstitials to be absorbed by\ndislocation loops. Thus, the average loop size decreases, as shown in Fig. 2c.\nCorresponding changes have been made in the revised manuscript (P6 Lines 152-154).\nFig. R1. Variations of DLs density and size with radiation dose at room temperature in\nCG and GNS samples. This is an enlarged view of Fig. 2c showing the situation at low\ndoses.\n4. P6 Line 160: What reactor system is T91 currently being used in?\nREPLY: T91 steel is widely used as structural material in nuclear reactors mostly\nbecause of its high mechanical properties, good phase stability, reduced activity, and\nhigh resistance to stress corrosion cracking. For example, in current fission reactors,\nT91 is used to manufacture pressure vessels, piping systems, and power boilers [10-12].\nIn the sodium-cooled and lead-cooled fast fission reactors, T91 has also been selected\nas a candidate for fuel cladding, duct/wrapper, and structural materials [13, 14].\nFurthermore, it is also an important candidate for structural components in future fusion\npower reactors [15] and concentrated solar power plants [16]. And T91 steel has also\nbeen considered as a candidate structural material in accelerator-driven subcritical\nsystem (ADS) served for nuclear reactors [17].\nThe applications of T91 steel in current reactor systems have been mentioned in the\nrevised manuscript (see P6 lines 172-173).\n5. P6: Line 167: Any comment on by which mechanism the voids are suppressed in GNS\ncomponents?\nREPLY: In GNS components, two mechanisms could accelerate the annihilation of\nradiation-induced displacement damages (mostly vacancies, interstitials, and their\nsmall clusters), and therefore suppress the formation of voids (recombined vacancies\nand small vacancy cluster) during the radiation process at high temperatures. The 1st\none is the absorption of vacancies and small vacancy clusters by a high density of grain\nboundaries, dislocation networks and stacking faults. Compared to state in normal\nnanocrystal samples, the density of above mentioned structures (grain boundaries,\ndislocation networks and stacking faults) in GNS is much higher, which could result in\nhigh probability for them to interact with their surrounding point defects (e.g. interstitial,\nvacancies and small defect clusters) induced by displacement cascade. Thus, the total\nnumber of left vacancies and small vacancy clusters should be lower than that in normal\nnanocrystal samples and CG samples. Furthermore, at low temperatures (e.g. less than\n300 ℃), diffusion of vacancies and their clusters is limited, thus, the formation of voids\nthrough absorption of these left vacancy defects has been suppressed. It should be noted\nthat while the microstructure remains stable after irradiation at 450 C, this mechanism\nstill works in GNS samples. The 2nd one to suppress formation of void is related to the\nAMT process explored in this work, which would annihilate the vacancies and small\nvacancy clusters by lattice rearrangement during the AMT process, as revealed by the\npresent MD simulation results (Supplementary Video 4). Therefore, the total number of\nvacancies and small vacancy clusters in both austenite and the transformed martensite\nto form a void is further decreased by AMT process.\nThe mechanisms by which the radiation resistance is enhanced in GNS components\nhave been analyzed in detail in Section Anti-radiation mechanisms. And some revisions\nhave been made to make the analysis more accurate and clearer in the revised\nmanuscript (P6 Line 177; P7 Lines 201-203; P8 Lines 235-237).\n6. P6 Line 171-174: These sentences are in contradiction to one another. Is refined\nmicrostructure an enhancement or an Achilles heel to radiation resistance? Radiation\nresistance should be understood in the context of defect sinks, while refined\nmicrostructure propensity for grain growth can be understood in the context of stored\nenergy within the microstructure.\nREPLY: Sorry that these sentences are not precise and bring misunderstanding. In fact,\nthe refined microstructure provides a high density of grain boundaries, which absorbs\nradiation defects and significantly increases the radiation resistance. However, this\nmechanism can only work well when the grain size is kept in the nanometer scale. When\na large number of radiation defects has been absorbed by nanostructured materials\nunder irradiation, the diffusion of grain boundaries would occur, which results in\nOstwald ripening process and finally leads to the instability of refined microstructure\nas reported previously [18, 19]. At this stage, these materials are not with refined\nmicrostructures and thus don’t show significantly enhanced radiation resistance\nanymore.\nCorresponding corrections have been made to avoid misunderstandings in the revised\nmanuscript. As shown in P7 Lines 184-189, the sentences have been changed to “One\nshould be its specific nanostructures with a high stability. Conventionally, refining\nmicrostructure leads to an enhanced radiation resistance, mostly due to the existence of\nnumerous GBs acting as sinks of radiation induced defects. However, the instability of\nrefined microstructure is an Achilles heel of traditional NC materials under irradiation,\nmostly due to the high stored energy of GBs in them”.\n7. P7 Line 201: By what mechanism does the martensitic transformation occur directly\nfrom irradiation? What is the driving force in the absence of a mechanical load to\ninitiate the transformation?\nREPLY: According to the results of high-resolution STEM observations and MD\nsimulations, we inferred that the martensitic transformation occurs mostly due to the\ninteractions between pre-existing SFs and radiation-induced defects in irradiated GNS\n304NG. As shown in Figs. 3(b, c) and Supplementary Fig. 21(c, d) in the revised\nmanuscript, the martensitic phase transition occurs mainly at the junctions of SFs after\nirradiation, which conforms to the classical Kurdjumov-Sachs relationship. And further\nMD simulations indicated that the phase transition is initiated while the radiation-\ninduced defects are absorbed by the SFs (Supplementary Video 4), confirming that the\nradiation induced phase transformation is firstly induced by the interaction between SFs\nand radiation-induced defects. And the mechanism of large-scale AMT was also\nanalyzed by building a MD model (Supplementary Fig. 25a) with cross SFs around the\nnewly formed martensite phase. Under the interaction of SFs and radiation-induced\ndefects, the phenomenon that martensite extends along these interactions to both sides\n(Supplementary Fig. 25b, Supplementary Video 5) is confirmed. By analyzing the\ndisplacement of interstitial atoms, it is found that the displacement of defects near the\ninterface between SFs and martensite is more significant, and finally martensite occurs\nat the intersection of SFs and extends to both sides (Supplementary Fig. 25c).\nMeanwhile, the corresponding phase transformation also promotes atomic\ndisplacements after the formation of new martensite phase, resulting in continuous\nmartensitic transformation and related defect absorption from the damage peak to the\nsurrounding regions.\nIt is noted that the application of external stress is just to accelerate the phase\ntransformation process without affecting the underlying mechanism during MD\nsimulations in this work, although inhomogeneous stress fields do exist in irradiated\nsamples induced by implantation of energetic particles [20]. For example, a previous\nstudy [20] demonstrated that the ion implantation could result in concentration gradient\nand stress fields in the near-surface layer, and the maximum value could reach several\nhundred MPa. Furthermore, displacement cascades and the formation of radiation\ndefects would also induce inhomogeneous stress fields around these regions [21].\nCorrespondingly, the driving force of martensitic transformation in irradiated GNS\nsamples mostly comes from the excess free energies relating with the pre-existing SFs\nand radiation-induced defects. Meanwhile, the residual stresses introduced by SMRT,\nphase transition, chemical segregation under irradiation and energetic particle\nirradiation, might also provide additional driving force. And their contributions have\nbeen discussed in the manuscript (see P10 Lines 273-277, Lines 302-307).\nSome other changes have also been made to further clarify the AMT mechanism in the\nrevised manuscript (see P10 Lines 273-274, Lines 280-284, and Supplementary Note\n1).\n8. P8 Line 218: What does adaptive martensitic transformation mean, and how is this\ndifferent from a standard martensitic transformation?\nREPLY: According to experimental and simulative results in this work, we found that\nthe martensitic transformation in irradiated GNS 304NG samples occurs much faster\nand in a broader region than that a standard martensitic transformation in irradiated\ncoarse-grained counterparts. Therefore, we proposed the adaptive martensitic\ntransformation (AMT) to name the martensitic transformation process in GNS samples\nunder irradiation.\nIn detail, the AMT mechanism shows different features from the standard martensitic\ntransformation: 1) The AMT completes in a much lower dose in the peak damage region\nof GNS samples than a standard martensitic transformation under irradiation at room\ntemperature without the application of external stress. The AMT results in a full\ndistribution of martensite in the peak damage region after irradiation at 12 dpa; while a\nlarge fraction of austenite retains in the peak damage region after a standard martensitic\ntransformation even at 155 dpa. 2) The AMT progresses adaptively from the peak\ndamage region to the surrounding regions; while a standard martensitic transformation\noccurs only in the peak damage region. 3) The AMT can occur at high temperatures\n(e.g. 450 C) when no standard martensitic transformation can occur.\nCorresponding changes have been made in the revised manuscript (P8 Line 221; Lines\n240-243).\n9. P9 Line 247: These stacking faults group together to form an hcp metastable\nmartensitic phase, which should be observable in the sample.\nREPLY: Yes. HCP metastable martensitic phase (-martensite) can be formed by\noverlapping bundles of stacking faults on {111} planes in principle, so that -martensite\nshould be observed in irradiated samples. However, we haven’t detected the formation\nof -martensite in irradiated samples by TEM observations and XRD analyses in this\nwork.\nFollowing the comment of the referee, we carefully checked the FFT pattern in Fig. 3c\nin the manuscript (also shown in Fig. R2 here). It should be emphasized that this is an\nimage observed at the 1 nm scale in a grain. At the atomic scale (i.e. HRTEM or\nHRSTEM), although it is possible that the misarrangement of continuous two or more\natomic layers (as marked by orange dash lines in Fig. R2a) might introduce the\ndiffraction spots of HCP phase in corresponding FFT patterns in FCC metals, more\nanalyses are required to confirm this possibility since the FFT patterns only reflect the\nperiodic changes in atomic lattice arrangement. In this work, it can be determined that\nthe misarrangement layers are SFs rather than HCP metastable phases according to the\ndiffraction pattern theory, as explained below: Firstly, the diffraction spots of γ and α'-\nmartensite can be easily identified (marked by red and green, respectively, in Fig. R2b).\nSecondly, according to the lattice constants of ε-martensite (a = b = 2.53 Å, c = 4.12\nÅ), a spot which seems to match the (101̅0)𝜀 should be observed (circled in orange\nin Fig. R2b). However, the D of the spot is 8.15 nm-1, while the calculated value of\n𝐷 = 9.13 nm-1 (error > 10%). The possibility that this diffraction spot is from -\n(101̅0)𝜀\nmartensite should be excluded. Therefore, the bundles should still be SFs, and the\nformation of an HCP phase by overlapping SFs bundles, within a thickness less than\n0.5 nm, should be excluded in this figure.\nHowever, further investigations are expected to reveal the formation of metastable -\nmartensite in irradiated samples, especially with the contribution from SFs.\nThis point has been explained in the revised manuscript (P10 Lines 284-287).\nFig. R2. (also Fig. 3c in the manuscript) The transformation in the lattice arrangement\nfrom  to  phase at SFs in GNS 304NG under irradiation at 2 dpa. a, A high-resolution\nSTEM image. b, The FFT pattern of (a) with marks showing the diffraction spots of ,\n’, and expected  phases.\n10. P9 Line 258: If this were to be the case, then once the MT was complete you would\nexpect to see an increase in the DL size and density\nREPLY: Yes. The DL size and density should increase after the completion of MT.\nAccording to the proposed AMT mechanism in this work, the radiation resistance\nshould be significantly enhanced due to lattice rearrangement during the AMT. When\nthe AMT is not complete, the surrounding regions under AMT will be strong sinks for\nradiation-induced defects, which should result in a lower DL size and density. When\nthe AMT is complete, no such a mechanism processes in the radiation affected range\nanymore, so that the DL size and density is expected to increase. It is noted that the\nformed BCC martensitic phase should be more radiation resistant than the original FCC\naustenitic phase, so that there should be a balance between an enhancing radiation\nresistance (due to increasing martensite fraction) and a decreasing resistance (due to\nreducing AMT ability) with an increasing radiation dose. As demonstrated in\nSupplementary Fig. 15 in the revised manuscript, in comparison with the DL density in\nGNS 304 counterparts possessing a similar grain size distribution but much more pre-\nexisting martensite (near 100%), an enhanced irradiation resistance can only be\nobserved in GNS 304NG samples at peak doses > 6 dpa, when a large fraction of\nmartensite has been formed and the AMT mechanism is still active. No increase in the\nDL size and density was observed in irradiated GNS 304NG in Fig. 2c possibly due to\nthat the AMT is not complete even after irradiation at a peak dose of 155 dpa (see Fig.\n3a) in this work.\nCorresponding modifications have been made in the revised manuscript (P10 Lines\n290-294).\n11. P26 Line 761: The bundles of SF here are likely hcp martensite formations. Examine\nthis described in\nhttps://www.sciencedirect.com/science/article/abs/pii/S0921509323014533\nhttps://www.sciencedirect.com/science/article/abs/pii/S1359645415003559\nREPLY: Yes. HCP martensite might be formed by overlapping bundles of SFs on {111}\nplanes in austenitic stainless steels under severe plastic deformation in principle as\ndescribed in these references, and the bundles of SFs show a similar morphology to\nthose of hcp martensite in Supplementary Fig. 5. However, when we checked the\nmicrostructure here, we found the bundles are still FCC austenitic phase instead of HCP\nmartensitic phase, as revealed in the SAED pattern in Fig. R3. Of course, the\nappearance of HCP martensite can’t be excluded completely in SMRT samples even we\ndidn’t observe such a phase by TEM and XRD in this work.\nThis point has been explained in the revised manuscript (Figure caption of\nSupplementary Fig. 5).\nFig. R3. (A part of Extended Data Figure 2 in the manuscript) Microstructure of GNS\n304NG at a depth of ~4 μm. a, The cross-sectional BF TEM image. b, The\ncorresponding SAED pattern.\nReference\n[1] Sun, C. et al. Superior radiation-resistant nanoengineered austenitic 304L stainless\nsteel for applications in extreme radiation environments. Sci. Rep. 5, 7801 (2015).\n[2] Hoffman, A. et al. Enhanced Resistance to Irradiation Induced Ferritic\nTransformation in Nanostructured Austenitic Steels. Materialia 13, 100806 (2020).\n[3] Sun, Y. T., Kong, X. & Wang, Z. B. Superior mechanical properties and deformation\nmechanisms of a 304 stainless steel plate with gradient nanostructure. Int. J. Plast. 155,\n103336 (2022).\n[4] Cheng, Z., Zhou, H. F., Lu, Q. H., Gao, H. J. & Lu, L. Extra strengthening and work\nhardening in gradient nanotwinned metals. Science 362, 559 (2018).\n[5] Li, X., Lu, L., Li, J., Zhang, X. & Gao, H. Mechanical properties and deformation\nmechanisms of gradient nanostructured metals and alloys. Nat. Rev. Mater. 5, 706-723\n(2020).\n[6] Ru, X. & Staehle, R. W. Historical Experience Providing Bases for Predicting\nCorrosion and Stress Corrosion in Emerging Supercritical Water Nuclear Technology.\nCorrosion 69, 211–229 (2013).\n[7] Barabash, V. et al. Materials challenges for ITER – Current status and future\nactivities. J. Nucl. Mater. 367–370, 21–32 (2007).\n[8] Gong, X., Short, M. P., Auger, T., Charalampopoulou, E. & Lambrinou, K.\nEnvironmental degradation of structural materials in liquid lead- and lead-bismuth\neutectic-cooled reactors. Prog. Mater. Sci. 126, 100920 (2022).\n[9] Osetsky, Y. N., Bacon, D. J., Serra, A., Singh, B. N. & Golubov, S. I. Stability and\nmobility of defect clusters and dislocation loops in metals. J. Nucl. Mater. 276, 65-77\n(2000).\n[10] Swindeman, R. W., Santella, M. L., Maziasz, P. J., Roberts, B. W. & Coleman, K.\nIssues in replacing Cr–Mo steels and stainless steels with 9Cr–1Mo–V steel. Int. J.\nPress. Vessels Pip. 81, 507-512 (2004).\n[11] Klueh, R. L. Elevated temperature ferritic and martensitic steels and their\napplication to future nuclear reactors. Int. Mater. Rev. 50, 287-310 (2013).\n[12] Shang, Z. et al. Gradient nanostructured steel with superior tensile plasticity. Sci.\nAdv. 9, eadd9780 (2023).\n[13] Grégoire, B., Oskay, C., Meißner, T. M. & Galetz, M. C. Corrosion mechanisms\nof ferritic-martensitic P91 steel and Inconel 600 nickel-based alloy in molten chlorides.\nPart II: NaCl-KCl-MgCl2 ternary system. Sol. Energy Mater. Sol. Cells 216, 110675\n(2020).\n[14] Thomas Davis. Review of the Iron-Based Materials Applicable for the Fuel and\nCore of Future Sodium Fast Reactors (SFR). (2018).\n[15] Davis, T. P. et al. Nanocluster evolution and mechanical properties of ion irradiated\nT91 ferritic-martensitic steel. J. Nucl. Mater. 548, 152842 (2021).\n[16] Yeli, G., Strutt, V. C. I., Auger, M. A., Bagot, P. A. J. & Moody, M. P.\nCharacterisation of nano-scale precipitates in BOR60 irradiated T91 steel using atom\nprobe tomography. J. Nucl. Mater. 543, 152466 (2021).\n[17] Vogt, J. B. et al. Understanding the liquid metal assisted damage sources in the\nT91 martensitic steel for safer use of ADS. Eng. Fail. Anal. 14, 1185-1193 (2007).\n[18] Wang, J., Xue, Z. & Song, S. Austenite Grain Boundary Precipitation and Growth\nBehavior of AlN Inclusions During Heat Treatment of Low-Density Steel: Ostwald\nRipening. Metall. Mater. Trans. B 56, 3271-3277 (2025).\n[19] Lescoat, M.L. et al. Radiation-induced Ostwald ripening in oxide dispersion\nstrengthened ferritic steels irradiated at high ion dose. Acta Mater. 78, 328–340 (2014).\n[20] Wang, Q., Ishikawa, H., Nakano, S., Ogiso, H. & Akedo, J. Curvature-based\nresidual stress measurement for ion-implanted stainless-steel sheet. Vacuum 75, 225-\n229 (2004).\n[21] Marcos, V. L., Asemann, P., Schrempel, F., Ronning, C. & Wendler, E. Wendler. Defect\ninduced stress in ion irradiated nanocrystalline Ge2Sb2Te5. Mater. Lett. 319, 132249\n(2022).\nReviewer #2:\nThe draft “Extraordinary Radiation Resistance in Metals via Adaptive Martensitic\nTransformation” by authors presents the findings of an experimental investigation\ndemonstrating the formation of unexpectedly lower densities of dislocation loops and\nvoids in a designed nanocrystalline austenitic stainless steel under ion irradiation. \nAccordingly, they find that a novel and interesting mechanism called adaptive\nmartensitic transformation (AMT) contributes to the superior radiation resistance. And\nthey further clarify and confirm the findings using MD simulations. Meanwhile, they\nalso find superior mechanical properties in the irradiated nanocrystalline samples.\nThe experimental design is reasonable. The microscopy studies are thorough and the\nmechanical measurements are credible. And the process to produce the designed\nnanocrystalline structure for AMT in mm-scale thicknesses at low cost is desirable. In\nthe referee’s opinion, this draft can be accepted for publication after several issues be\naddressed.\nREPLY：Thank you for the positive comments and insightful suggestions.\n1. Martensitic transformation is a topic worthy of in-depth research, such as TRIP effect,\nshape memory alloys, magnetism, etc. Residual stress has a significant impact on the\nmartensitic transformation. Although the authors mentioned the release of residual\nstress at high temperatures in the line 261-267, the effect of residual stress on the AMT\nmechanism at room temperature deserves further investigation and may become a\nmeans of surface modification. Do the authors pay attention to the relationship between\nmartensitic transformation and residual stress in the irradiation samples?\nREPLY: Yes. We have analyzed the relationship between martensitic transformation\nand residual stress in GNS samples under irradiation in this work.\nResidual stress might play an important role in martensitic transformation in austenitic\nsteels, TRIP steels, shape memory alloys, et al. Typically, compressive stresses promote\nmartensitic transformation by decreasing the energy barrier of transformation. Since\nhigher compressive residual stresses exist in SMRT 304NG samples than in CG\nsamples, the AMT process might be further promoted in addition to the contribution of\nSFs in GNS samples under irradiation at room temperature. In this work, the role of\nresidual stress on martensitic transformation has also been revealed by MD simulations,\nin which remarkable stresses were introduced so that the martensitic transformation was\naccelerated in an austenitic unit (see Supplementary Fig. 21). Meanwhile, it is noticed\nthat the large-scale AMT and enhanced irradiation resistance have also been observed\nin GNS sample irradiated at higher temperatures (see Supplementary Fig. 10), when the\nresidual stresses are almost completely released. Therefore, it should be SFs play the\ndominant role in the AMT process in this work. However, further investigations are\nexpected to clarify the effect of residual stress on the AMT under irradiation at room\ntemperature.\nMeanwhile, we have measured the residual stresses in the surface layer of GNS and CG\nsamples before and after irradiation at room temperature. As shown in Fig. R4, a certain\namount of compressive stress was introduced into the surface layer of 304NG samples\nby SMRT. As the AMT progresses, the residual compressive stress was further\nincreased (from -131.215.4 MPa to -182.610.7 MPa), while there is no detectable\nchange in the residual stress in irradiated CG counterparts.\nIn addition, as pointed by the reviewer, martensitic transformation is also an issue\nrelating with TRIP (transformation induced plasticity) effect, and therefore radiation\nmight bring significant changes in TRIP materials. For example, Agrawal et al. [1]\nfound that radiation might initiate a self-healing mechanism in TRIP high entropy alloys\n(HEAs) by balancing radiation induced martensitic transformation with temperature\ninduced reverse transformation, so that a high stability of mechanical properties is\nachieved in HEAs under irradiation at high doses and temperatures. And the self-\nhealing mechanism might be introduced to expand the service dose and temperature of\nTRIP HEAs in future reactors.\nCorresponding modifications have been made in the revised manuscript (P10 Lines\n302-307 on residual stress; P2 Lines 65-69, P9 Line 268 and P13 Line 376 on TRIP\neffect).\nFig. R4. The variation of residual stress with irradiation dose of CG and GNS 304NG\nsample.\n2. The manuscript claims that the AMT mechanism is related to high-density SFs. So,\nwhat is the value of it. A detailed value would be appreciated.\nREPLY: Yes. We have checked the density of SFs in detail. The derived value is\n~7.5×1014 m-2. The density of SFs has been added in the revised manuscript (P4 Line\n111).\n3. The authors focus on studying the microstructure of the peak damage region.\nAccording to the dose and ion distribution map provided in Supplementary Fig. 3,\nthe irradiation dose and ion concentration in the peak damage area are much higher\nthan other irradiation regions. So, will this affect the martensitic transformation in this work?\nREPLY: As presented in this work, the proposed AMT process can in fact extend from\nthe peak damage region to the surrounding regions until the phase transformation is\ncompletely finished. It should be noted that the surrounding regions around the peak\ndamage region have different irradiation doses and implanted ion concentrations.\nHowever, the AMT also occurs in these regions. Therefore, it can be known that the\ndifferent irradiation doses and ion concentrations have no or slight effects on the\noccurrence of AMT process. It should also be noted that the ions used in this work is\nFe-ions, which become Fe atoms and may not affect the formation of DLs and voids in\nirradiated 304NG.\nCorresponding changes have been made in the revised manuscript (P8 Lines 240-243).\n4. In lines 105-106, the authors mention a 3D dislocation network was formed in\n304NG by SMRT. That’s interesting. However, there seems to be no experimental\nevidence, and only been demonstrated through MD simulations.\nREPLY: Although the TEM images are 2D projected, they can also reflect a certain of\nspatial information. For example, as shown in Fig. R5, some white lines (i.e. dislocation\nlines, indicated by blue dashed arrows) are observed in the nano-grains and can pass\nthrough the interface, and when numerous dislocation lines intertwining with each other,\na dislocation network has been formed in the whole GNS surface region. This\ndislocation network structure is thus distributed in both grains and interfaces, forming\na special dislocation structure in GNS materials, as demonstrated in previous studies\n[2]. Furthermore, such a DN structure is essentially a type of GNDs, because it is\ngenerated to accommodate strain gradient induced by microstructure gradients.\nThe dislocation lines which forming DNs have been marked in Fig. 1d and\ncorresponding changes (see P4 Line 113) have been made in the revised manuscript.\nFig. R5. High magnification STEM image of GNS 304NG sample. The yellow solid\narrow indicate SFs and blue dashed arrows indicate DNs.\n5. In lines 141-143, “it is interesting to notice that both the density and size of DLs\nsignificantly decrease……0.5 to 5 dpa……” the range of the dose is incorrect (refer to\nSupplementary Fig. 4).\nREPLY: Thank referee for pointing this out! Following your suggestion, we have\ncorrected it (2 to 10 dpa) and included the correct information in the revised manuscript.\n(see P6 Line 151; P7 Line 208)\nReference\n[1] Agrawal, P. et al. Irradiation response of innovatively engineered metastable TRIP\nhigh entropy alloy. J. Nucl. Mater. 574, 154217 (2023).\n[2] Cheng, Z., Zhou, H. F., Lu, Q. H., Gao, H. J. & Lu, L. Extra strengthening and\nwork hardening in gradient nanotwinned metals. Science 362, 559 (2018).\nReviewer #3:\nThis manuscript states the radiation tolerance of metals can be enhanced by a\nmechanism called adaptive martensitic transformation. The manuscript uses the\nexperimental methods and molecular dynamics simulations to reveal this mechanism.\nAfter reviewing this manuscript, the reviewer has some comments. See details in below.\nREPLY：Thank you for the thorough reviewing and constructive comments.\n1. Is the first paragraph the abstract? If this is an abstract, we should not cite the\nreferences. Please remove these references and reorganize the references in the whole\nmanuscript.\nREPLY: Yes. This paragraph is the abstract. Some references were provided here since\nthis submission was directly transferred from another journal which requires articles\nstart with a referenced summary.\nCorrespondingly, we have removed the references and reorganized the references in the\nrevised manuscript.\n2.This manuscript is comparing the radiation behaviors of the CG and GNS material,\nI would suggest revealing the grain size distribution for both CG and GNS. The\nmanuscript only listed the average grain size of CG is 26 µm, what is the grain size and\nits distribution for GNS?\nREPLY: As suggested by the reviewer, we have added the arithmetic mean grain sizes\nand corresponding uncertainties, 2610 m for CG material and 3713 nm for GNS\nmaterial, in the revised manuscript (P14 Line 386 and P4 Line 107, respectively).\n3.This manuscript discussed the phase transformation under irradiations. What about\nthe radiation induced grain growth, especially under the high dose after 155 dpa? The\nauthors should also reveal the grain size after 155 dpa and explain why the material is\nstill within nanometer scale after high irradiation dose in more details. I can see the\nauthors mentioned combining low energy interfaces, SFs, and DNs contributes to a\nhigh stability from both thermodynamic and kinetic perspectives. Please explain more\nhow these interfaces, SFs and DNs will contribute the structure stability in\nthermodynamic and kinetic perspectives.\nREPLY: The microstructure is rather stable in GNS 304NG samples under irradiation,\nas shown in Supplementary Fig. 6 in this work. Statistical measurements revealed that\nthe arithmetic mean grain size and uncertainty are 8626 nm after irradiation 155 dpa\nat room temperature (Fig. R6) and are 7232 nm after irradiation 80 dpa at 450 °C.\nThe contributions of low energy interfaces (e.g. Σ3 (111) GBs), SFs and DNs to the\nstructure stability can be understood as follows: At first, low energy interfaces are\ntypically very stable under heating or irradiation mostly due to their low stored energies\n[1, 2]. In this work, a massive density of Σ3 (111) GBs are formed in GNS samples,\nwhich contribute to the high microstructure stability from a thermodynamic aspect.\nSecondly, SFs, especially the junctions of SFs, could induce the AMT and significantly\nannihilate the radiation induced defects inside grains. This should decrease the defects\nabsorption at GBs and therefore reduce GB migration under irradiation. Therefore, the\nexistence of a massive pre-formed SFs also contributes to microstructure stability.\nFinally, DNs would not only absorb radiation induced defects inside grains and reduce\nGB migration, but also interact with GBs and obstacle their migration by forming\ninterfacial Lomer-Cottrell locks, junctions, and so on [3]. Therefore, such a special\nmicrostructure (combining low energy interfaces, SFs, and DNs) contributes to a high\nstability from both thermodynamic and kinetic perspectives. And the grain size is still\nwithin nanometer scale after irradiation at a very high dose or at a high temperature.\nCorresponding explanations have been added in the revised manuscript (P7 Lines 189-\n203).\n35\n155dpa GNS 304NG\nArithmetic: 86.0 ± 25.9 nm\n25\n)\n%\n(\ny\n20\nc\nn\ne\nu15\nq\ne\nr\nF\n10\n5\n0\n20 40 60 80 100 120 140 160 180\nGrain size (nm)\nFig. R6. The grain size distribution of 155 dpa GNS 304NG.\n4. I suggest all the SAED patterns in this manuscript should be identified. Especially\nthe diffraction rings for the martensite in the GNS images.\nREPLY: Yes. All the SAED patterns have been indexed with more details in\nSupplementary Fig. 19 in the revised manuscript.\n5. In Figure 2, this manuscript is discussing the DL size of the CG and GNS sample.\nHowever, the error bar for the CG sample is way bigger than the GNS sample. Can you\nexplain this?\nREPLY: The arithmetic mean value and standard deviation of DL sizes are presented\nfor each sample in Fig. 2 and Supplementary Table 3. As exampled in Supplementary\nFig. 7, the distribution range of DL sizes in CG samples (2-15 nm) irradiated 155 dpa\nis larger than that in GNS counterparts (2-7 nm). This leads a bigger error bar in\nstatistical results of CG samples. And the underlying mechanism should be related with:\n1) much more radiation induced defects have been absorbed by grain boundaries,\ndislocation networks and stacking faults in GNS, resulting in smaller DLs. Further\nannihilation of point defects by AMT process would continuously reduce interstitials to\nDLs, thus, the error bar of DL size in GNS is expected to be small. 2) After AMT process,\nmartensitic phase has been formed, in which the diffusion activation energy of DLs is\nlow, resulting in strong absorption of DLs by grain boundaries. Therefore, less DLs are\nexpected to be observed, which would thus have smaller mean value and error bar in\nDL size.\nCorresponding modifications have been made in the revised manuscript (P6 Lines 145-\n149). And the mean values and uncertainties of DL sizes have been presented in\nSupplementary Fig. 7 and Supplementary Table 3.\n6.Based on the discussion and Figure 2, it looks like after 10 dpa, the DL size and\ndensity remain flat. Please explain why 10 dpa is the critical irradiation point in more\ndetails? Why will the DL size not increase, and DL density will not decrease after 10\ndpa? Is this related to the martensitic phase?\nREPLY: Yes. The fact that the DL size and density remain near flat after 10 dpa should\nbe related to the AMT process in GNS samples, while these values increase gradually\nwith an increasing dose in CG counterparts. As shown in Supplementary Fig. 6 and Fig.\n3a, at the initial stage of irradiation, the AMT process has not been triggered and the\ndefects formed in austenitic phase dominate the DL evolution process. With increase of\nradiation dose, further AMT process has been observed around the peak damage region\nto annihilate the point defects and small loops, resulting in decreases of both density\nand size of DLs. When the annihilation of point defects and small loops and the defect\nformation by irradiation in both martensite and surrounding austenite reach in an\nequilibrium state, both density and size of DLs also reach an equilibrium state until the\nphase transformation has been finished in the whole damage region. Thus, as the referee\nimplied, it is a process to reach the equilibrium state below the radiation dose of 12 dpa,\nwhich is related to the AMT process and the formed martensitic phase.\nSome details can also be found in the reply to Comments 8 and 10 from Reviewer #1.\nAccordingly, we have clarified the AMT mechanism with more details, by which the\nvariations of DL density and size with radiation dose should be more understandable in\nthe revised manuscript (P8 Lines 240-243).\n7. Please check all the references and keep them in the same format. Especially the\nformat for the journal’s name.\nREPLY: Yes. We have checked all references and corrected the them in the same format\nin the revised manuscript.\n8. In extended data figure 3, why did not put the images for 155 dpa? It only includes\nthe images up to 85 dpa. Since 155 dpa is the final dose, I would suggest adding the\nimages for 155 dpa.\nREPLY: In fact, the TEM image of the GNS sample irradiated at 155 dpa has been\nincluded in Extended Data Figure 3 in the last version (Supplementary Fig. 6 in the\nrevised manuscript). In Supplementary Fig. 6j, we can see that the martensitic\ntransformation has occurred around the irradiation region in the GNS sample irradiated\nunder an ion fluence of 1.625×1017 cm−2, i.e., a dose of 155 dpa in the peak damage\narea.\nFor a better understanding, we have rearranged the images in Supplementary Fig. 6 in\nthe revised manuscript.\n9. In extended data figure 4, the figure 4 (b) has no scale bar. Also, it looks larger of\nthe DL in Figure d than Figure c. Please clarify this.\nREPLY: A scale bar has been added in this figure (moved to Supplementary Fig. 7) in\nthe revised manuscript.\nThe reason why the DLs look larger in (d) than in (c) in this figure is as following:\nTypical TEM images reflect microstructure projected from a 3D space to a 2D space.\nIn the characterization of bright- and dark-field TEM images, the imaging contrasts\nmainly include mass-thickness contrast and diffraction contrast. A larger appearance in\n(d) was originated from the diffraction contrast in our last version TEM images. When\nthe diffraction conditions deviate from the Bragg diffraction conditions, the contrast of\nDLs will appear inside or outside them. And the deviation is related to g (diffraction\nvector), b (Burgers vector of dislocation loops), and s (deviation parameters of Bragg\ndiffraction conditions) according to the work of Föll and Wilkens [4]:\n(𝑔·𝑏)·𝑠 < 0,𝑜𝑢𝑡𝑠𝑖𝑑𝑒,\n(𝑔·𝑏)·𝑠 > 0,𝑖𝑛𝑠𝑖𝑑𝑒\nThe schematic illustration of inside-outside contrast is also shown in Fig. R7. To\nconfirm our clarification, by rotating a same grain to -g condition (Fig. R8), it can be\nseen that the diffraction contrast of identifiable DLs is inside the DLs, and visually,\ntheir size is significantly reduced compared to the previous image. It should be\nemphasized that the actual size of DLs will not change due to the influence of contrast\nimages. Therefore, when the size of a DL is measured, it should be based on its actual\nboundary. Of course, the images we provided previously did cause a significant\nmisunderstanding.\nThank you again for this suggestion! We have made the corrections in Supplementary\nFig. 7 in the revised manuscript.\nFig. R7. The schematic illustration of inside-outside contrast in TEM.\nFig. R8. The microstructure of GNS 304NG samples irradiated at room temperature for\n155 dpa. The BF TEM images of (a) g=011 and (b) -g, respectively. c, d, The\ncorresponding WBDF TEM images.\n10. In extended data figure 6, it says the white arrows show g= [001]. However, it is\nvery difficult to see in the images, please improve the image quality.\nREPLY: The image quality of this figure (moved to Supplementary Fig. 12) has been\nimproved in the revised manuscript.\n11. In the supplementary Figure 3, I would suggest changing the curve colors in the\nright image to make it more obvious to see the difference.\nREPLY: The curves have been noted by different symbols to make the difference more\nobvious in the right image of this figure (moved to Supplementary Fig. 4) in the revised\nmanuscript.\nReference\n[1] Lu, K. Stabilizing nanostructures in metals using grain and twin boundary\narchitectures. Nat. Rev. Mater. 1, 16019 (2016).\n[2] Zhang, X. et al. Radiation damage in nanostructured materials. Prog. Mater. Sci. 96,\n217–321 (2018).\n[3] Zhu, Y., Li, Z. & Huang, M. Atomistic modeling of the interaction between matrix\ndislocation and interfacial misfit dislocation networks in Ni-based single crystal\nsuperalloy. Comput. Mater. Sci. 70, 178–186 (2013).\n[4] Föll, H. & Wilkens, M. A simple method for the analysis of dislocation loops by\nmeans of the inside-outside contrast on transmission electron micrographs. Phys. Status\nSolidi A 31, 519-524 (1975)."}, {"round": 2, "reviewer_comments": {"reviewer_1": "All comments have been adequately addressed.", "reviewer_2": "The scientific content of this manuscript is interesting and holds certain scientific value. Therefore, I believe that this revised manuscript is ready for publication.", "reviewer_3": "The authors have addressed all my comments. I think this manuscript is ready for publish."}, "author_response": ""}]

Decision letter