Available via license: CC BY 4.0
Content may be subject to copyright.
Citation: He, W.; Liu, F.; Tan, L.; Tian,
Z.; Qin, Z.; Huang, L.; Xiao, X.; Wang,
G.; Chen, P.; Liu, B. Optimizing the
Thermomechanical Process of
Nickel-Based ODS Superalloys by an
Efficient Method. Materials 2022,15,
4087. https://doi.org/10.3390/
ma15124087
Academic Editors: Tomasz
Trzepieci´nski and Valentin ¸Stefan
Oleksik
Received: 24 April 2022
Accepted: 26 May 2022
Published: 9 June 2022
Publisher’s Note: MDPI stays neutral
with regard to jurisdictional claims in
published maps and institutional affil-
iations.
Copyright: © 2022 by the authors.
Licensee MDPI, Basel, Switzerland.
This article is an open access article
distributed under the terms and
conditions of the Creative Commons
Attribution (CC BY) license (https://
creativecommons.org/licenses/by/
4.0/).
materials
Article
Optimizing the Thermomechanical Process of Nickel-Based
ODS Superalloys by an Efficient Method
Wuqiang He 1,2, Feng Liu 1,2 , Liming Tan 1,2, Zhihui Tian 3, Zijun Qin 1,2, Lan Huang 1,2, Xiangyou Xiao 1,2,
Guowei Wang 1,2, Pan Chen 4,* and Baogang Liu 5,*
1State Key Laboratory of Powder Metallurgy, Central South University, Changsha 410083, China;
183301019@csu.edu.cn (W.H.); liufeng@csu.edu.cn (F.L.); limingtan@csu.edu.cn (L.T.);
zijun.qin@csu.edu.cn (Z.Q.); lhuang@csu.edu.cn (L.H.); xiangyou.xiao@csu.edu.cn (X.X.);
wangguowei23@csu.edu.cn (G.W.)
2Research Institute of Powder Metallurgy, Central South University, Changsha 410083, China
3College of Science, Huazhong Agriculture University, Wuhan 430070, China;
2017301210207@webmail.hzau.edu.cn
4School of Minerals Processing and Bioengineering, Central South University, Changsha 410083, China
5School of Energy and Electromechanical Engineering, Hunan University of Humanities, Science and
Technology, Loudi 417000, China
*Correspondence: panchen@csu.edu.cn (P.C.); liudd2016@126.com (B.L.); Tel.: +86-0731-8883-0938 (P.C.);
+86-0731-8883-0938 (B.L.)
Abstract:
Thermo-mechanical process of nickel-based oxide dispersion strengthened (ODS) superal-
loys is critical to produce desired components. In this study, an efficient method of consolidating
powder is introduced to optimize the preparation process, microstructure and properties of nickel-
based ODS superalloys. The influences of consolidation temperature, strain rate and ball milling
time on the hardness of nickel-based superalloys were studied. The relationship among process,
microstructure and hardness was established, the nanoparticles strengthening and grain boundary
strengthening in nickel-based ODS superalloys were discussed. The results indicate that long ball
milling time, moderately low consolidation temperature and high strain rates are beneficial to im-
proving properties of nickel-based superalloys. Moreover, dispersion strengthening of nanoparticles
and grain boundary strengthening play important roles in enhancing nickel-based ODS superalloys.
Keywords: nickel-based ODS superalloys; consolidation; strengthening; hardness
1. Introduction
Nickel-based oxide dispersion strengthened (ODS) superalloys have been considered
as promising candidate materials such as gas turbines, heat exchanger tubes and nuclear
reactors due to their excellent radiation resistance, comprehensive mechanical proper-
ties and high temperature creep strength [
1
–
5
]. These enhanced properties are mainly
attributed to high number density nanoscale particles, which can pin grain boundaries
and dislocations [
6
–
8
]. The addition such as YH
2
and Y
2
O
3
is dissociated to Y during
the mechanical alloying (MA) process, then Y reacts with and Al or Ti in the matrix to
form nanoscale Y-Al-rich or Y-Ti-rich oxides (e.g., YAlO
3
, Y
3
Al
5
O
12
,Y
4
Al
2
O
9
, Y
2
Ti
2
O
7
,
Y
2
TiO
5
) in consolidation process [
9
]. The hot extrusion (HEX) and hot isostatic pressing
(HIP) are commonly used consolidation processes to prepare nickel-based ODS superalloys.
However, the preparation of ball milling and consolidation of powder usually takes several
months and thus it is costly, which limits the development of ODS superalloys.
It is well known that the properties of nickel-based superalloys including nickel-based
ODS superalloys are strongly dependent on the processing parameters. For example, Kim
et al. [
10
] investigated the influence of several processing parameters including milling tem-
perature, rotation speed, and consolidation temperature on microstructure and properties
of 14Cr ODS steels alloy. As milling temperature decreases, milling speed increases, and
Materials 2022,15, 4087. https://doi.org/10.3390/ma15124087 https://www.mdpi.com/journal/materials
Materials 2022,15, 4087 2 of 25
HIP temperature decreases, the microstructures of samples become finer and more uniform,
and considerable improvement of tensile strength was noticed. Kishimoto et al. [
11
] and
Odette et al. [
12
,
13
] reported that with the increase of HIP temperature the grain size of the
alloy grew significantly, and the dislocation density decreased, the average particle size
of oxides increased and the density of oxides decreased. Tan et al. [
14
] studied the effects
of HIP temperature and pressure on the grain in PM superalloy FGH96, and it was found
that temperature and pressure played different roles in controlling prior particle boundary
(PPB) precipitation and grain structure during HIP, the tendency of grain coarsening under
high temperature could be inhibited by increasing HIP pressure which facilitated the re-
crystallization. It is essential to optimize the processing parameters to obtain components
with desired microstructure and properties. Therefore, to accelerate the development of
nickel-based superalloys, some efficient experimental methods need to be developed [15].
In this study, a new efficient thermal consolidation method was used to prepare nickel-
based superalloys with and without ODS. It takes only a few minutes and a few grams of
powder to prepare a sample by this method, which dramatically shortens the preparation
cycle and reduces the cost. In addition, the influences of consolidation temperature, strain
rate and ball milling time on the hardness of nickel-based superalloys were studied, and the
relationship of process, microstructure and properties was established, the strengthening
mechanism of nickel-based superalloys with and without ODS were discussed.
2. Materials and Experiment Methods
The chemical composition of two nickel-based powders (PA1 and PA2) prepared
by argon atomization in this work was given in Table 1. MA powder was prepared by
ball milling PA powders and 0–0.6 wt.% YH
2
powders for 0–36 h with speed 350 rpm in
high-energy planetary mill filled with high purity argon gas (99.999%). The ball-to-powder
mass ratio was 8:1 and 0.5 wt.% ethanol was added as process control agent. The MIX
powder was prepared by mixing MA and PA powder at ratio 1:2 (wt.%) for 12 h with speed
60 rpm. The laser particle size analyzer (Mastersizer 3000, Malvern Panalytical, Malvern,
UK) was adopted to detected the particle size distribution of powder.
Table 1. Chemical compositions of PA1 and PA2 powder in wt.%.
Al Cr Fe Ti Y C Ni
PA1 0.25 20.5 0.67 0.56 0 0.054 Bal.
PA2 0.25 21 0.85 0.57 0.68 0.059 Bal.
The powder was sealed in stainless steel container with 10 mm diameter and 15 mm
height, vacuum electron beam welding is adopted to welded container under a vacuum
of 10
−2
torr. Thereafter, the container was compressed by Gleeble 3180D at temperatures
ranging from 850
◦
C to 1150
◦
C with strain rates 0.1 s
−1
–5 s
−1
. Carbon sheets and lubricant
were placed between specimen and dies to make the deformation more uniform during
compression. The specimens were originally heated to certain temperatures with a rate
of 5
◦
C/s and held for 3 min to homogenize temperature of specimen. To monitor the
temperature during the compression, the thermal-couples were welded at the longitudinal
center of the specimen surface. Finally, the specimen was compressed according to the pre-
set procedure and immediately water quenched to freeze the deformation microstructure.
All specimens were compressed to a total engineering strain of 80% with constant strain
rates. The process flow diagram of specimens prepared is indicated in Figure 1, and the
processing methods of specimens have been compared in Table 2.
Materials 2022,15, 4087 3 of 25
Materials2022,15,xFORPEERREVIEW3of25
Figure1.Flow‐processdiagramofsamplespreparation.
Table2.Comparisonofprocessingmethodsofspecimens.
SpecimenPowderMillTime/hTemperature/°CStrainRate/s
−1
PA1PA1powder0
850/950/
1050/11500.1/1/5
MA1‐24hPA1powderand0.6%YH
2
powder24
MA1‐36hPA1powderand0.6%YH
2
powder36
MIX1‐36hMA1‐36hpowderandPA2powder(ratio1:2)12(mixed)
PA2PA2powder0
MA2‐24hPA2powder24
MA2‐36hPA2powder36
MIX2‐24hMA2‐24hpowderandPA2(ratio1:2)12(mixed)
ThephasesofpowderandalloyswereinvestigatedusingX‐raydiffractometer(XRD,
D/MAX2500,RIGAKU,Tokyo,Japan)withCuKαradiationof0.15405nm.Thediffrac‐
tionAngle(θ)rangesfrom20°to80°withscanningrateof1°/min.
TheVickersmicrohardnessmeasurementswereperformedbyusingvickershard‐
nesstester(THV‐10,TEST‐TECH,Shanghai,China)ataloadof3000gandadwelltime
of10s.Thestandarddeviationwascalculatedbasedon10measurements.
Thecharacterizationofmicrostructurealloyswasobservedbyopticalmicroscopy
(OM,DM4000M,LEICA,Germany)andfield‐emissionSEM(Quanta650,FEI,Hillsboro,
OR,USA)equippedwithanelectronbackscatterdiffraction(EBSD),thecentralofsamples
wasusedtoanalyzed,asshowninFigure1.Thespecimensweremechanicallypolished
byabrasivepapersand50nmaluminiumoxide,thespecimensforOMobservationwere
etchedbyasolutionof100mLethanol+100mLHCl+5gCuCl
2
,andthespecimensfor
EBSDobservationwerevibrationpolishedfor3–8h.Thevibrationpolishingwascon‐
ductedonBuehlervibratorypolisherVibroMet2(Buehler,USA)with30%amplitude.The
EBSDwasconductedat20kVacceleratingvoltage,0.05or1μmstepsize,thedatawas
Figure 1. Flow-process diagram of samples preparation.
Table 2. Comparison of processing methods of specimens.
Specimen Powder Mill Time/h Temperature/◦C Strain Rate/s−1
PA1 PA1 powder 0
850/950/
1050/1150 0.1/1/5
MA1-24h PA1 powder and 0.6%YH2powder 24
MA1-36h PA1 powder and 0.6%YH2powder 36
MIX1-36h MA1-36h powder and PA2 powder (ratio 1:2) 12 (mixed)
PA2 PA2 powder 0
MA2-24h PA2 powder 24
MA2-36h PA2 powder 36
MIX2-24h MA2-24h powder and PA2 (ratio 1:2) 12 (mixed)
The phases of powder and alloys were investigated using X-ray diffractometer (XRD,
D/MAX 2500, RIGAKU, Tokyo, Japan) with Cu K
α
radiation of 0.15405 nm. The diffraction
Angle (θ) ranges from 20◦to 80◦with scanning rate of 1 ◦/min.
The Vickers microhardness measurements were performed by using vickers hardness
tester (THV-10, TEST-TECH, Shanghai, China) at a load of 3000 g and a dwell time of 10 s.
The standard deviation was calculated based on 10 measurements.
The characterization of microstructure alloys was observed by optical microscopy
(OM, DM4000M, LEICA, Germany) and field-emission SEM (Quanta 650, FEI, Hillsboro,
OR, USA) equipped with an electron backscatter diffraction (EBSD), the central of samples
was used to analyzed, as shown in Figure 1. The specimens were mechanically polished
by abrasive papers and 50 nm aluminium oxide, the specimens for OM observation were
etched by a solution of 100 mL ethanol + 100 mL HCl + 5 g CuCl
2
, and the specimens
for EBSD observation were vibration polished for 3–8 h. The vibration polishing was
conducted on Buehler vibratory polisher VibroMet 2 (Buehler, USA) with 30% amplitude.
The EBSD was conducted at 20 kV accelerating voltage, 0.05 or 1
µ
m step size, the data was
analyzed via HKL Channel 5 software and the equivalent grain size was determined based
on grain area A, 2√A/π.
Materials 2022,15, 4087 4 of 25
Transmission electron microscope (TEM, Themis Z 3.2, FEI, USA) observation with
energy dispersive spectrometer (EDS) was used to characterize the structure of specimens,
under 200 kV accelerating voltage. The TEM samples were polished to a thickness of
50
µ
m by abrasive papers, and sectioned to diameter 3mm. Then the slices were twin-jet
electropolished at
−
25
◦
C and 40 V in the corrosive solution of 90% ethanol and 10%
perchloric acid.
3. Results and Discussion
3.1. Morphologies of Powder Analysis
Figure 2a,c show the SEM images of precursor powders, it is evident that PA powder
is spherical or subspherical, the EBSD inverse pole figure (IPF), grain distribution and grain
misorientation are shown in Figure 2b–f. The grain size ranges from 0.5
µ
m to 18
µ
m, the
average grain size of PA1 and PA2 powder is 4.8
µ
m and 5.8
µ
m, and the average grain
misorientation of PA1 and PA2 powder is 33.8 and 31.4, respectively.
The microstructure and particle size distribution of powder with different ball milling
time are shown in Figure 3. The particle size of PA powder ranges in 1–170
µ
m, the
powder particle is irregular morphology after ball milling, because of the repeated and
high-speed collision between powder and powder or powder and ball milling medium
during MA. Large plastic deformation of powder after crushing during the ball milling
process. Within ranges of 20–600
µ
m, the particle size increases with the increasing ball
milling time, specifically, the average particle sizes of PA1 and PA2 powder are 182
µ
m
and 224
µ
m after ball milling 24 h, and these of PA1 and PA2 powder increase to 205
µ
m
and 245
µ
m after 36 h ball milling. The average particles size of powder increases due to
welding of powders during ball milling.
3.2. Macroscopic Cracking Analysis
96 specimens were prepared at temperatures ranging from 850
◦
C to 1150
◦
C with
strain rates of 0.1–5 s
−1
. The shapes of the specimens after hot compression are shown
in Figure 4. From macroscopic perspective, most of the specimens are regular pie with
obvious bulging in the middle, which is due to the non-uniform deformation caused by the
friction between the specimen and workpiece during hot compression [
16
]. Some specimens
with obviously macroscopic failure features (marked by the red box) are observed, these
specimens are compressed at the low temperature, the cracking percentage of the samples
is 14.6% and 5.2% compressed at 850
◦
C and 950
◦
C, respectively. No macroscopic cracking
was observed in samples consolidated at 1050 ◦C and 1150 ◦C.
Figure 5depicts some typical features of PA1 and MA1-36h alloys consolidated at
different temperatures with strain rate of 0.1 s
−1
. The microscopic cracks along the PPB
are observed in PA1 alloys and MA1-36h alloys at 850
◦
C corresponding to Figure 5a,c,
respectively. Generally, PPB are difficult to be broken and eliminated at low temperature,
which facilitates the crack nucleation and propagation during thermal consolidation [
17
].
The microcrack along the PPB are mainly caused by local stress concentration [
18
,
19
].
With the increase of consolidation temperature, the thermal activation of the material
increases, the atomic diffusion rate accelerates, and the kinetic energy of the atoms increases,
weakening the binding force between atoms, the softening effect of the material becomes
obvious. The PPB is broken and eliminated quickly at elevated temperature and high stress,
replaced by dynamic recrystallized grains, and good metallurgical bonding between the
powder particles is achieved [
20
]. No PPB and macroscopic cracks are observed in PA1 and
MA1-36h alloy consolidated at 1050 ◦C.
Materials 2022,15, 4087 5 of 25
Materials2022,15,xFORPEERREVIEW5of25
Figure2.TheSEMandEBSDIPFimagesofpowders:(a,b)PA1;(c,d)PA2;(e)grainsizedistribution;
(f)grainmisorientation.
Figure 2.
The SEM and EBSD IPF images of powders: (
a
,
b
) PA1; (
c
,
d
) PA2; (
e
) grain size distribution;
(f) grain misorientation.
Materials 2022,15, 4087 6 of 25
Materials2022,15,xFORPEERREVIEW6of25
Figure3.Themicrostructuralfeatureandparticlesizedistributionfor(a)PA1,(b)MA1‐24h,(c)
MA1‐36h,(d)MIX1‐36h,(e)PA2,(f)MA2‐24h,(g)MA2‐36h,(h)MIX2‐24h,respectively.
3.2.MacroscopicCrackingAnalysis
96specimenswerepreparedattemperaturesrangingfrom850°Cto1150°Cwith
strainratesof0.1–5s
−1
.Theshapesofthespecimensafterhotcompressionareshownin
Figure4.Frommacroscopicperspective,mostofthespecimensareregularpiewithobvi‐
ousbulginginthemiddle,whichisduetothenon‐uniformdeformationcausedbythe
frictionbetweenthespecimenandworkpieceduringhotcompression[16].Somespeci‐
menswithobviouslymacroscopicfailurefeatures(markedbytheredbox)areobserved,
thesespecimensarecompressedatthelowtemperature,thecrackingpercentageofthe
samplesis14.6%and5.2%compressedat850°Cand950°C,respectively.Nomacroscopic
crackingwasobservedinsamplesconsolidatedat1050°Cand1150°C.
Figure 3.
The microstructural feature and particle size distribution for (
a
) PA1, (
b
) MA1-24h, (
c
) MA1-
36h, (d) MIX1-36h, (e) PA2, (f) MA2-24h, (g) MA2-36h, (h) MIX2-24h, respectively.
Materials 2022,15, 4087 7 of 25
Materials2022,15,xFORPEERREVIEW7of25
Figure4.Theshapesofspecimenscompressedatdifferentconditions:(a)PA1;(b)PA2;(c)MA1‐24
h;(d)MA2‐24h;(e)MA1‐36h;(f)MA2‐36h;(g)MIX1‐36h;(h)MIX2‐24h.
Figure5depictssometypicalfeaturesofPA1andMA1‐36halloysconsolidatedat
differenttemperatureswithstrainrateof0.1s
−1
.ThemicroscopiccracksalongthePPBare
observedinPA1alloysandMA1‐36halloysat850°CcorrespondingtoFigure5a,c,re‐
spectively.Generally,PPBaredifficulttobebrokenandeliminatedatlowtemperature,
whichfacilitatesthecracknucleationandpropagationduringthermalconsolidation[17].
ThemicrocrackalongthePPBaremainlycausedbylocalstressconcentration[18,19].With
theincreaseofconsolidationtemperature,thethermalactivationofthematerialincreases,
theatomicdiffusionrateaccelerates,andthekineticenergyoftheatomsincreases,weak‐
eningthebindingforcebetweenatoms,thesofteningeffectofthematerialbecomesobvi‐
ous.ThePPBisbrokenandeliminatedquicklyatelevatedtemperatureandhighstress,
replacedbydynamicrecrystallizedgrains,andgoodmetallurgicalbondingbetweenthe
powderparticlesisachieved[20].NoPPBandmacroscopiccracksareobservedinPA1
andMA1‐36halloyconsolidatedat1050°C.
Figure 4.
The shapes of specimens compressed at different conditions: (
a
) PA1; (
b
) PA2; (
c
) MA1-24 h;
(d) MA2-24 h; (e) MA1-36 h; (f) MA2-36 h; (g) MIX1-36 h; (h) MIX2-24 h.
Materials 2022,15, 4087 8 of 25
Materials2022,15,xFORPEERREVIEW8of25
Figure5.TheSEMimagesoffailurespecimensconsolidatedatdifferenttemperatureswithstrain
rate0.1s
−1
:(a)PA1:850°C;(b)PA1:1050°C;(c)MA1‐36h:850°C;(d)MA1‐36h:1050°C.
Theformationofcrackinhotcompressionaredescribedbythethreestepsinthe
schematic,asshowninFigure6.Atinitialstages,loosepowderbegantoslideandrear‐
rangementunderlowpressure.Thereafter,asthestressincrease,theplasticdeformation
ofpowdergraduallyhappens,whenthestressreachestheyieldpointofthepowder.As
hotcompresscontinue,strainincreasesgradually,dynamicrecrystallization(DRX)oc‐
curs.However,whenhotcompressiontemperatureislow,PPBcanhardlybebroken,and
graduallyisstretchedintoellipsoidasstrainincreases[20],leadingtotheconcentration
ofstressatPPB,microcracksgeneratesandexpandsalongthePPB,leadingtothefor‐
mationofmacroscopiccrackseventually.
Figure 5.
The SEM images of failure specimens consolidated at different temperatures with strain
rate 0.1 s−1: (a) PA1: 850 ◦C; (b) PA1: 1050 ◦C; (c) MA1-36h: 850 ◦C; (d) MA1-36h: 1050 ◦C.
The formation of crack in hot compression are described by the three steps in the
schematic, as shown in Figure 6. At initial stages, loose powder began to slide and rear-
rangement under low pressure. Thereafter, as the stress increase, the plastic deformation
of powder gradually happens, when the stress reaches the yield point of the powder. As
hot compress continue, strain increases gradually, dynamic recrystallization (DRX) occurs.
However, when hot compression temperature is low, PPB can hardly be broken, and gradu-
ally is stretched into ellipsoid as strain increases [
20
], leading to the concentration of stress
at PPB, microcracks generates and expands along the PPB, leading to the formation of
macroscopic cracks eventually.
Materials 2022,15, 4087 9 of 25
Materials2022,15,xFORPEERREVIEW9of25
Figure6.Schematicillustrationofformationofcompressioncracksduringhotcompression.(I)prior
todeformation,(II)deformation,(III)theformationofcracks.
3.3.XRDofthePowderandAlloysAnalysis
Figure7ashowstheXRDpatternsofprecursorandballmillingpowder.γ‐Niand
YH
2
diffractionpeaksarevisibleinthemixturepowderofPAandYH
2
powder.Afterball
milling36h,onlyγ‐NidiffractionpeakareobservedinMApowder,YH
2
isunstableand
dissociatedtoYandHduringtheMA,whichisconsistentwithotherstudies[21].Mean‐
while,theimpactenergygeneratedbytheballmillingresultsinareductionofgrainsize
andanaccumulationofthelatticestrain,whichleadstothedecreaseofdiffractionpeak
intensity,increaseofwidthandangulardeviation[22].TheXRDpatternsofPA1and
MA1‐36halloyconsolidatedat1050°C/0.1s
−1
isillustratedinFigure7b,thediffraction
peaksofY
4
Al
2
O
9
andTiO
2
aredetectedinMA1‐36halloy,onlythediffractionofTiO
2
can
befoundinPA1alloy.
Figure7.XRDpatternsof(a)powdersmilledfordifferenttimesand(b)samplesconsolidatedat
1050°C/0.1s
−1
.
3.4.MicrostructureCharacterization
ThegrainsatcenterregionsofPA1andMA1‐36halloyconsolidatedatdifferentcon‐
ditionsarepresentedbyEBSDIPF,asshowninFigures8and9.Generally,thesamples
consolidatedathighertemperaturesandlowerstrainrateshavelargeaveragegrainsize.
Significantgraingrowthhappenswhenthetemperatureexcesses950°C,especiallyatlow
strainrate.Astheconsolidationtemperatureincreasesfrom850°Cto1150°Catstrain
rateof0.1s
−1
,theaveragegrainsizeofPA1alloyincreasesfrom2.76μmto16.5μm,and
thatoftheMA1‐36halloyincreasesfrom0.18μmto0.31μm.Thegrainevolutionisgen‐
erallycontrolledbydynamicrecovery(DRV),DRX,andgraingrowthprocesses[23–26].
Figure 6.
Schematic illustration of formation of compression cracks during hot compression. (
I
) prior
to deformation, (II) deformation, (III) the formation of cracks.
3.3. XRD of the Powder and Alloys Analysis
Figure 7a shows the XRD patterns of precursor and ball milling powder.
γ
-Ni and
YH
2
diffraction peaks are visible in the mixture powder of PA and YH
2
powder. After
ball milling 36 h, only
γ
-Ni diffraction peak are observed in MA powder, YH
2
is unstable
and dissociated to Y and H during the MA, which is consistent with other studies [
21
].
Meanwhile, the impact energy generated by the ball milling results in a reduction of grain
size and an accumulation of the lattice strain, which leads to the decrease of diffraction
peak intensity, increase of width and angular deviation [
22
]. The XRD patterns of PA1 and
MA1-36h alloy consolidated at 1050
◦
C/0.1 s
−1
is illustrated in Figure 7b, the diffraction
peaks of Y
4
Al
2
O
9
and TiO
2
are detected in MA1-36h alloy, only the diffraction of TiO
2
can
be found in PA1 alloy.
Materials2022,15,xFORPEERREVIEW9of25
Figure6.Schematicillustrationofformationofcompressioncracksduringhotcompression.(I)prior
todeformation,(II)deformation,(III)theformationofcracks.
3.3.XRDofthePowderandAlloysAnalysis
Figure7ashowstheXRDpatternsofprecursorandballmillingpowder.γ‐Niand
YH
2
diffractionpeaksarevisibleinthemixturepowderofPAandYH
2
powder.Afterball
milling36h,onlyγ‐NidiffractionpeakareobservedinMApowder,YH
2
isunstableand
dissociatedtoYandHduringtheMA,whichisconsistentwithotherstudies[21].Mean‐
while,theimpactenergygeneratedbytheballmillingresultsinareductionofgrainsize
andanaccumulationofthelatticestrain,whichleadstothedecreaseofdiffractionpeak
intensity,increaseofwidthandangulardeviation[22].TheXRDpatternsofPA1and
MA1‐36halloyconsolidatedat1050°C/0.1s
−1
isillustratedinFigure7b,thediffraction
peaksofY
4
Al
2
O
9
andTiO
2
aredetectedinMA1‐36halloy,onlythediffractionofTiO
2
can
befoundinPA1alloy.
Figure7.XRDpatternsof(a)powdersmilledfordifferenttimesand(b)samplesconsolidatedat
1050°C/0.1s
−1
.
3.4.MicrostructureCharacterization
ThegrainsatcenterregionsofPA1andMA1‐36halloyconsolidatedatdifferentcon‐
ditionsarepresentedbyEBSDIPF,asshowninFigures8and9.Generally,thesamples
consolidatedathighertemperaturesandlowerstrainrateshavelargeaveragegrainsize.
Significantgraingrowthhappenswhenthetemperatureexcesses950°C,especiallyatlow
strainrate.Astheconsolidationtemperatureincreasesfrom850°Cto1150°Catstrain
rateof0.1s
−1
,theaveragegrainsizeofPA1alloyincreasesfrom2.76μmto16.5μm,and
thatoftheMA1‐36halloyincreasesfrom0.18μmto0.31μm.Thegrainevolutionisgen‐
erallycontrolledbydynamicrecovery(DRV),DRX,andgraingrowthprocesses[23–26].
Figure 7.
XRD patterns of (
a
) powders milled for different times and (
b
) samples consolidated at
1050 ◦C/0.1 s−1.
3.4. Microstructure Characterization
The grains at center regions of PA1 and MA1-36h alloy consolidated at different
conditions are presented by EBSD IPF, as shown in Figures 8and 9. Generally, the samples
consolidated at higher temperatures and lower strain rates have large average grain size.
Significant grain growth happens when the temperature excesses 950
◦
C, especially at low
strain rate. As the consolidation temperature increases from 850
◦
C to 1150
◦
C at strain rate
of 0.1 s
−1
, the average grain size of PA1 alloy increases from 2.76
µ
m to 16.5
µ
m, and that
of the MA1-36h alloy increases from 0.18
µ
m to 0.31
µ
m. The grain evolution is generally
Materials 2022,15, 4087 10 of 25
controlled by dynamic recovery (DRV), DRX, and grain growth processes
[23–26]
. For the
PA1 alloy, comparing with the initial microstructure of precursor powder of PA1 (Figure 2b),
grains are not uniform and some large grains only deform under stress rates of 0.1 s
−1
and
1 s
−1
at 850
◦
C. Grains of alloy consolidated at 950
◦
C are more uniform than that of the
samples consolidated at 850
◦
C, with the increase of thermal activation. Grain growth plays
a dominating role during the microstructure evolution when the temperature excesses
950
◦
C. Comparing with microstructure of the PA1 alloys, the grains of MA1-36h alloys
consolidated at 850–950
◦
C keep fine during consolidation, and the average grain size
fluctuates slightly. The grains grow gradually in samples as the consolidation temperature
increasing to 1050
◦
C. Moreover, it is clear that the grains decrease from micron to submicron
by comparing PA1 alloy with MA1-36 alloy, the grains are refined by ball milling. The grain
growth rate of PA1 alloy is more rapidly than that of MA1-36h alloy with the temperature
above 1050
◦
C, mechanism behind that will be discussed thereafter. In terms of strain
rate, the strain rate has ambiguous effects on the grain evolution at 850–950
◦
C, as the
consolidation temperature excesses 1050
◦
C, the average grain size of samples consolidated
at 5 s
−1
is smaller than that of other strain rates, and at a certain temperature, grains in
alloys deformed at 5 s
−1
are more uniform than these of the counterparts consolidated at
other strain rates. Accordingly, the consolidation temperature of near 1050
◦
C with strain
rate of 5 s−1is suggested, to obtain homogeneous distributed fine-grains.
Figure 10 shows the grain misorientations of PA1 and MA1-36h alloys consolidated
at different conditions. The influence of consolidation conditions on the grain evolution
can be reflected by grain boundary misorientation [
27
]. The dislocation occurs in the
process of consolidation, which leads to the misorientation in grains, and forming low
angle grain boundaries (LAGBs, <15
◦
). The higher the dislocation density, the larger the
proportion of LAGBs. DRX is triggered, when the dislocation density reaches critical
dislocation density. The dislocation is absorbed during the grain growth, resulting in the
decrease of dislocation density, the fraction of LAGBs decreases. With sufficient driving
force and time, DRX continues to occur and the dislocation density decreases further. For
the PA1 alloy, comparing with PA1 powder (Figure 2), the average grain misorientation of
samples decreases, as shown in Figure 10a. Generally, As the consolidation temperature
increases from 850
◦
C to 1050
◦
C, the thermal activation increases, the fraction of high
angle grain boundary (HAGBs, >15
◦
) and the average grain misorientation of samples
increases. However, the average grain misorientation decreases at 1150
◦
C, it is speculated
that the grain growth leads to the decrease of the grain quantity in the same area. For the
MA1-36h alloy, the average grain misorientation of samples increases with the increase of
temperature except 1150 ◦C/1 s−1(Figure 10b).
In addition, Figure 11 shows the TEM bright field images of PA1 and MA1-36h alloy
consolidated at 1150
◦
C/5 s
−1
. Figure 11a shows a large-grain region with a high dislocation
density in the PA1 alloy. The dislocations have long, relatively straight-line segments, with
sharp corners and a serrated appearance in the other locations. Figure 11b illustrates the
microstructure of MA1-36h alloy, significant nanoparticles and dislocations are noticed in
sample, the dislocations interact with the nanoparticles, indicating that they are strongly
pinned by nanoparticles. The nanoparticles are distributed in grain and grain boundary,
ranging in size from a few nanometers to tens of nanometers.
Materials 2022,15, 4087 11 of 25
Materials2022,15,xFORPEERREVIEW11of25
Figure 8. Cont.
Materials 2022,15, 4087 12 of 25
Materials2022,15,xFORPEERREVIEW12of25
Figure8.EBSDmapsatcenterregionof(a)PA1and(b)MA1‐36halloysconsolidatedatdifferent
conditions.
Figure 8.
EBSD maps at center region of (
a
) PA1 and (
b
) MA1-36h alloys consolidated at different
conditions.
Materials 2022,15, 4087 13 of 25
Materials2022,15,xFORPEERREVIEW13of25
Figure9.Theequivalentgrainsizedistributionsandaveragegrainsizesof(a,c)PA1and(b,d)MA1‐
36halloysconsolidatedatdifferentconditions.
Figure10showsthegrainmisorientationsofPA1andMA1‐36halloysconsolidated
atdifferentconditions.Theinfluenceofconsolidationconditionsonthegrainevolution
canbereflectedbygrainboundarymisorientation[27].Thedislocationoccursinthepro‐
cessofconsolidation,whichleadstothemisorientationingrains,andforminglowangle
grainboundaries(LAGBs,<15°).Thehigherthedislocationdensity,thelargerthepropor‐
tionofLAGBs.DRXistriggered,whenthedislocationdensityreachescriticaldislocation
density.Thedislocationisabsorbedduringthegraingrowth,resultinginthedecreaseof
dislocationdensity,thefractionofLAGBsdecreases.Withsufficientdrivingforceand
time,DRXcontinuestooccurandthedislocationdensitydecreasesfurther.ForthePA1
alloy,comparingwithPA1powder(Figure2),theaveragegrainmisorientationofsam‐
plesdecreases,asshowninFigure10a.Generally,Astheconsolidationtemperaturein‐
creasesfrom850°Cto1050°C,thethermalactivationincreases,thefractionofhighangle
grainboundary(HAGBs,>15°)andtheaveragegrainmisorientationofsamplesincreases.
However,theaveragegrainmisorientationdecreasesat1150°C,itisspeculatedthatthe
graingrowthleadstothedecreaseofthegrainquantityinthesamearea.FortheMA1‐
36halloy,theaveragegrainmisorientationofsamplesincreaseswiththeincreaseoftem‐
peratureexcept1150°C/1s
−1
(Figure10b).
Figure 9.
The equivalent grain size distributions and average grain sizes of (
a
,
c
) PA1 and
(b,d) MA1-36h alloys consolidated at different conditions.
Materials 2022,15, 4087 14 of 25
Materials2022,15,xFORPEERREVIEW14of25
Figure 10. Cont.
Materials 2022,15, 4087 15 of 25
Materials2022,15,xFORPEERREVIEW15of25
Figure10.Thegrainmisorientationsof(a)PA1and(b)MA1‐36halloysconsolidatedatdifferent
conditions.
Inaddition,Figure11showstheTEMbrightfieldimagesofPA1andMA1‐36halloy
consolidatedat1150°C/5s
−1
.Figure11ashowsalarge‐grainregionwithahighdislocation
densityinthePA1alloy.Thedislocationshavelong,relativelystraight‐linesegments,
withsharpcornersandaserratedappearanceintheotherlocations.Figure11billustrates
themicrostructureofMA1‐36halloy,significantnanoparticlesanddislocationsareno‐
ticedinsample,thedislocationsinteractwiththenanoparticles,indicatingthattheyare
stronglypinnedbynanoparticles.Thenanoparticlesaredistributedingrainandgrain
boundary,ranginginsizefromafewnanometerstotensofnanometers.
ThegraingrowthrateofMA1‐36halloyisretardedcomparingwithPA1alloy.Since
nanoparticlesdetectedasY
4
Al
2
O
9
byXRDaredispersedinMA1‐36halloy,whichprevent
Figure 10.
The grain misorientations of (
a
) PA1 and (
b
) MA1-36h alloys consolidated at different
conditions.
Materials 2022,15, 4087 16 of 25
Materials2022,15,xFORPEERREVIEW16of25
DRVandgraingrowthviapinningdislocationsandgrainboundaries[28],thegrain
boundariesofMA1‐36halloyaremorestablethanthatofPA1alloy,evenatveryhigh
temperature[29].
Figure11.TEMphotographsof(a)PA1and(b)MA1‐36halloycompressedat1150°C/5s
−1
.
ThelocalmisorientationimagesofPA1alloyandMA1‐36halloyconsolidatedatdif‐
ferentconditionsillustratethehardeningstatusofgrains,aspresentedinFigure12.In
general,thespecimensconsolidatedathighertemperatureshowslowerlocalmisorienta‐
tion.ThestrainhardeningofPA1consolidatedat850°Cissignificant,asshowninFigure
12a,whichisrelatedtodislocationaccumulationinmaterials.Withtheincreaseoftem‐
perature,thestrainhardeninggetsrelievedbyDRVathighertemperature,whichalso
facilitatestheDRXviasubgrainformation[30].Inaddition,asthestrainrateincreases
from0.1s
−1
to5s
−1
,thePA1alloypresentslowerlocalmisorientationexceptat1050°C.
However,thestrainratehasambiguouseffectsonthelocalmisorientationevolutionin
MA1‐36halloy,asshowninFigure12b.
EBSDmapsinFigures13and14depictthedeformed,substructuredandrecrystal‐
lizedgrainsinthespecimensconsolidatedatdifferentconditions.Ingeneral,withthe
increaseofconsolidationtemperature,thefrequencyofdeformedgrainsdeducesgradu‐
ally,whilefractionsoftherecrystallizedgrainsgetlargeratthesamestrainrates.Ingen‐
eral,thedeformedgrainsofPA1alloydecreasesgraduallywithconsolidationstrainrate
increasingfrom0.1s
−1
to5s
−1
.Astonishingly,thesubstructuredgrainsinMA1‐36halloy
consolidatedatstrainratesof1s
−1
aremorethanthecounterpartsconsolidatedatstrain
ratesof0.1s
−1
and5s
−1
.
Figure 11. TEM photographs of (a) PA1 and (b) MA1-36h alloy compressed at 1150 ◦C/5 s−1.
The grain growth rate of MA1-36h alloy is retarded comparing with PA1 alloy. Since
nanoparticles detected as Y
4
Al
2
O
9
by XRD are dispersed in MA1-36h alloy, which prevent
DRV and grain growth via pinning dislocations and grain boundaries [
28
], the grain
boundaries of MA1-36 h alloy are more stable than that of PA1 alloy, even at very high
temperature [29].
The local misorientation images of PA1alloy and MA1-36h alloy consolidated at differ-
ent conditions illustrate the hardening status of grains, as presented in Figure 12. In general,
the specimens consolidated at higher temperature shows lower local misorientation. The
strain hardening of PA1 consolidated at 850
◦
C is significant, as shown in Figure 12a, which
is related to dislocation accumulation in materials. With the increase of temperature, the
strain hardening gets relieved by DRV at higher temperature, which also facilitates the
DRX via subgrain formation [
30
]. In addition, as the strain rate increases from 0.1 s
−1
to
5 s
−1
, the PA1 alloy presents lower local misorientation except at 1050
◦
C. However, the
strain rate has ambiguous effects on the local misorientation evolution in MA1-36h alloy, as
shown in Figure 12b.
EBSD maps in Figures 13 and 14 depict the deformed, substructured and recrystallized
grains in the specimens consolidated at different conditions. In general, with the increase
of consolidation temperature, the frequency of deformed grains deduces gradually, while
fractions of the recrystallized grains get larger at the same strain rates. In general, the
deformed grains of PA1 alloy decreases gradually with consolidation strain rate increasing
from 0.1 s
−1
to 5 s
−1
. Astonishingly, the substructured grains in MA1-36h alloy consolidated
at strain rates of 1 s
−1
are more than the counterparts consolidated at strain rates of 0.1 s
−1
and 5 s−1.
Materials 2022,15, 4087 17 of 25
Materials2022,15,xFORPEERREVIEW17of25
Figure 12. Cont.
Materials 2022,15, 4087 18 of 25
Materials2022,15,xFORPEERREVIEW18of25
Figure12.EBSDmapsshowinglocalmisorientationingrainsof(a)PA1and(b)MA1‐36halloys
consolidatedatdifferentconditions,whereindifferentcolorsrangingfrom0°to5°correspondto
differentmisorientationsasindicatedbythecolorbar.
Figure 12.
EBSD maps showing local misorientation in grains of (
a
) PA1 and (
b
) MA1-36h alloys
consolidated at different conditions, wherein different colors ranging from 0
◦
to 5
◦
correspond to
different misorientations as indicated by the color bar.
Materials 2022,15, 4087 19 of 25
Materials2022,15,xFORPEERREVIEW19of25
Figure 13. Cont.
Materials 2022,15, 4087 20 of 25
Materials2022,15,xFORPEERREVIEW20of25
Figure13.EBSDimagesshowingthedistributionsofrecrystallized,substructured,anddeformed
grainsin(a)PA1and(b)MA1‐36halloysconsolidatedatdifferentconditions.
Figure 13.
EBSD images showing the distributions of recrystallized, substructured, and deformed
grains in (a) PA1 and (b) MA1-36h alloys consolidated at different conditions.
Materials 2022,15, 4087 21 of 25
Materials2022,15,xFORPEERREVIEW21of25
Figure14.Relativefrequencyofdeformed,substructuredandrecrystallizedgrainsin(a)PA1and
(b)MA1‐36halloysconsolidatedatdifferentconditions.
3.5.TheVariationofHardness
ThehardnessofspecimensmanufacturedatdifferentconditionsisshowninFigure
15.ThehardnessofPAalloyandMAalloyrangesin154.8–344.9HVand421.1–697.3HV,
respectively.Asexpected,thehardnessofMAalloyissignificantlyhigherthanthatofPA
alloy,andthehardnessincreaseswiththeballmillingtimeprolonging.Inaddition,the
hardnessofsamplesisalsostronglydependentupontemperature.Thehardnessofsam‐
plesdecreaseswiththeincreaseoftheconsolidationtemperature.Forexample,thehard‐
nessofPA1alloydecreasesfrom288.6HVto166.3HV,withtheconsolidationtempera‐
tureincreasesfrom850°Cto1150°Catstrainrateof5s
−1
.Theeffectofstrainrateonthe
alloyhardnessisalsopresented.Ingeneral,thehardnessofsamplesincreaseswithin‐
creasingstrainrate,especiallyat1050–1150°C.Highdeformationrateisconducivetothe
formationofhigh‐densitydislocation,higherenergystorageandrefinegrainsize,im‐
provingthepropertiesofthealloy,whichisconsistentwiththeotherliteraturereports
[31–38].
Figure 14.
Relative frequency of deformed, substructured and recrystallized grains in (
a
) PA1 and
(b) MA1-36h alloys consolidated at different conditions.
3.5. The Variation of Hardness
The hardness of specimens manufactured at different conditions is shown in Figure 15.
The hardness of PA alloy and MA alloy ranges in 154.8–344.9 HV and 421.1–697.3 HV,
respectively. As expected, the hardness of MA alloy is significantly higher than that of PA
alloy, and the hardness increases with the ball milling time prolonging. In addition, the
hardness of samples is also strongly dependent upon temperature. The hardness of samples
decreases with the increase of the consolidation temperature. For example, the hardness
of PA1 alloy decreases from 288.6 HV to 166.3 HV, with the consolidation temperature
increases from 850
◦
C to 1150
◦
C at strain rate of 5 s
−1
. The effect of strain rate on the alloy
hardness is also presented. In general, the hardness of samples increases with increasing
strain rate, especially at 1050–1150
◦
C. High deformation rate is conducive to the formation
of high-density dislocation, higher energy storage and refine grain size, improving the
properties of the alloy, which is consistent with the other literature reports [31–38].
Materials 2022,15, 4087 22 of 25
Materials2022,15,xFORPEERREVIEW22of25
Figure15.Thehardnessofalloyspreparedatdifferentballmillingtimeandconsolidationprocess.
4.Discussion
4.1.MicrostructureEvolutionduringConsolidation
ThegrainevolutionduringthermalconsolidationisgenerallycontrolledbyDRV,
DRXandgraingrowth[23–26].Thestrainhardeningissignificantinspecimenscom‐
pressedat850°C,adequateenergyisstoredinmaterialsbydislocationaccumulationto
triggerDRX,asillustratedbythelocalmisorientationmapsinFigure12.Astemperature
increasing,thestrainhardeninggetsrelieved,theimprovedtemperatureenhancesthe
formationofDRXnucleationviasubgrainformation,theinitiationofDRXisfasterat
highertemperature[30,39].Ingeneral,theimprovedtemperatureincreasestheaverage
grainsizeofsamplesbasedonFigure8andFigure9.Theimprovedtemperaturecanalso
enhancetheDRVandgraingrowthwhichreliefthedislocationaccumulationandcon‐
sumethestoredenergy,makingitlesssufficienttotriggersufficientrecrystallization.
Inaddition,intermsofstrainrate,thegraingrowthisrelativelyslackenedandaver‐
agegrainsizedropsduringconsolidationasthestrainrateincreasesfrom0.1s
−1
to5s
−1
,
speciallyat1050–1150°C.ThehigherstrainrateinhabitsDRVandlimitsthetimeof
boundarymigration,DRXnucleationviacoalescenceofsubgrainandthestraininduced
boundarymigrationisretarded[39,40].
4.2.StrengtheningMechanismofNickel‐BasedODSSuperalloys
ThehardnessH
V
ofsamplescanbedividedintothematrixhardnessH
0
,grainbound‐
arystrengtheningH
g
,solidsolutionstrengtheningH
ss
,oxidesstrengtheningH
p
anddislo‐
cationstrengtheningH
d
,itcanbeexpressedby[41]:
Figure 15. The hardness of alloys prepared at different ball milling time and consolidation process.
4. Discussion
4.1. Microstructure Evolution during Consolidation
The grain evolution during thermal consolidation is generally controlled by DRV, DRX
and grain growth [
23
–
26
]. The strain hardening is significant in specimens compressed at
850
◦
C, adequate energy is stored in materials by dislocation accumulation to trigger DRX,
as illustrated by the local misorientation maps in Figure 12. As temperature increasing,
the strain hardening gets relieved, the improved temperature enhances the formation of
DRX nucleation via subgrain formation, the initiation of DRX is faster at higher temper-
ature
[30,39]
. In general, the improved temperature increases the average grain size of
samples based on Figures 8and 9. The improved temperature can also enhance the DRV
and grain growth which relief the dislocation accumulation and consume the stored energy,
making it less sufficient to trigger sufficient recrystallization.
In addition, in terms of strain rate, the grain growth is relatively slackened and average
grain size drops during consolidation as the strain rate increases from 0.1 s
−1
to 5 s
−1
,
specially at 1050–1150
◦
C. The higher strain rate inhabits DRV and limits the time of
boundary migration, DRX nucleation via coalescence of subgrain and the strain induced
boundary migration is retarded [39,40].
4.2. Strengthening Mechanism of Nickel-Based ODS Superalloys
The hardness H
V
of samples can be divided into the matrix hardness H
0
, grain bound-
ary strengthening H
g
, solid solution strengthening H
ss
, oxides strengthening H
p
and dislo-
cation strengthening Hd, it can be expressed by [41]:
HV=H0+Hg+Hss +HP+Hd(1)
Materials 2022,15, 4087 23 of 25
In this work, the difference of PA alloy and MA alloy can be estimated by equation:
∆HV=∆Hg+∆HPs +∆HDis (2)
The contribution of dislocation strengthening is given by Equations (3) and (4) [
42
,
43
]:
∆Hd=αMGb√ρd(3)
ρd=2θ/µb(4)
where Mis Taylor factor equal to, Gis shear modulus, bis Burgers vector,
α
is the dislo-
cation strengthening coefficient,
ρd
is the dislocation density,
µ
is the unit length,
θ
is the
misorientation angle estimated by the local misorientation maps (Figure 12).
The grain boundary strengthening is related to the average grain size D, which can be
determined by the Hall-Petch equation [44]:
∆Hg=k
√D(5)
where kis the Halle-Petch strength constant. It is clear that the hardness of MA1-36
alloy is higher than that of PA1 alloy by comparing with the average grain size based on
Figures 8and 9.
In addition, nano-oxides plays an important role in ODS superalloy, as shown in
Figure 11. Based on Orowan strengthening mechanism [
45
,
46
], homogeneously dispersed
oxide particles can suppress effectively dislocation movement, as well as grain boundaries
migration [
47
,
48
]. Therefore, high hardness of MA1-36h alloy comes mainly from grain
boundary strengthening and nano-oxides strengthening.
5. Conclusions
In summary, nickel-based superalloys with and without ODS have been fabricated
by ball milling and efficient consolidation method. The relationship among the prepara-
tion process, microstructure and properties of alloy are studied. Basically, the following
conclusions can be reached:
(1)
The PPB of nickel-based superalloys is difficult to be broken and eliminated at low
temperature, which facilitates the crack nucleation and propagation during thermal
consolidation. With increasing of temperature, the thermal activation of the material
increases, the PPB is broken and eliminated quickly at elevated temperature and high
stress.
(2)
The grain size is sensitive to the consolidation temperature, the average grain size
increases with the increase of consolidation temperature. The average grain size of
samples consolidation at strain rate 5 s
−1
decreases and more uniform than low strain
rate, since higher strain rate inhabits DRV and limits the time of boundary migration.
(3)
The hardness of nickel-based superalloys decreases with the increase of the con-
solidation temperature, strain rate, and the hardness increases after ball-milled for
longer time. In addition, the hardness of nickel-based ODS superalloys is significantly
higher than that of nickel-based superalloys without ODS due to the grain boundary
strengthening and nano-oxides strengthening.
(4)
Basically, in order to obtain fine-grains, excellent properties and less cracking risk,
consolidation temperature of near 1050 ◦C and strain rate of 5 s−1are suggested.
Author Contributions:
Experimental design and project administration, W.H., F.L., L.T., L.H., B.L.,
P.C., Z.T., Z.Q., X.X. and G.W.; investigation and Formal analysis, Z.T., Z.Q., X.X., G.W., B.L. and
W.H.; writing—review & editing, W.H. and P.C. All authors have read and agreed to the published
version of the manuscript.
Materials 2022,15, 4087 24 of 25
Funding:
This work was supported by The National Science and Technology Major Project (Grant No.
2017-VI-0008-0078), The Natural Science Foundation of China (Grant No. 91860105, and 52074366),
and the Science and Technology Innovation Program of Hunan Province (2021RC3131). Lan Huang
acknowledges the Changsha Municipal Natural Science Foundation (kq2014126).
Institutional Review Board Statement: Not applicable.
Informed Consent Statement: Not applicable.
Data Availability Statement:
The data used to support the findings of this study are available from
the corresponding author upon request.
Acknowledgments:
This work was supported by the Project supported State Key Laboratory of
Powder Metallurgy, Central South University, Changsha, China.
Conflicts of Interest: The authors declare no conflict of interest.
References
1.
Verhiest, K.; Almazouzi, A.; Wispelaere, N.D.; Petrov, R.; Claessens, S. Development of oxides dispersion strengthened steels for
high temperature nuclear reactor applications. J. Nucl. Mater. 2009,385, 308–311. [CrossRef]
2.
De Carlan, Y.; Bechade, J.L.; Dubuisson, P.; Seran, J.L.; Billot, P.; Bougault, A.; Cozzikaa, T.; Doriota, S.; Hamona, D.; Henry, J.; et al.
CEA developments of new ferritic ODS alloys for nuclear applications. J. Nucl. Mater. 2009,386–388, 430–432. [CrossRef]
3.
Ukai, S.; Harada, M.; Okada, H.; Inoue, M.; Nomura, S.; Shikakura, S.; Nishidab, T.; Fujiwarab, M.; Asabe, K. Tube manufacturing
and mechanical properties of oxide dispersion strengthened ferritic steel. J. Nucl. Mater. 1993,204, 74–80. [CrossRef]
4. Benjamin, J.S.; Volin, T.E. The mechanism of mechanical alloying. Metall. Trans. 1974,5, 1929–1934. [CrossRef]
5.
Dong, H.; Yu, L.; Liu, Y.; Liu, C.; Li, H.; Wu, J. Effect of hafnium addition on the microstructure and tensile properties of aluminum
added high-Cr ODS steels. J. Alloy. Compd. 2017,702, 538–545. [CrossRef]
6.
Park, J.; Jang, J.; Kim, T.K.; Kim, S.-J.; Ahn, J.-H. Formation of Nanostructures in Ni-22Cr-11Fe-1X (X = Y
2
O
3
, TiO
2
) Alloys by
High-Energy Ball-Milling. J. Nanosci. Nanotechnol. 2011,11, 6213–6218. [CrossRef]
7.
Srolovitz, D.J.; Luton, M.J.; Petkovic-Luton, R.; Barnett, D.M.; Nix, W.D. Diffusionally modified dislocation-particle elastic
interactions. Acta Metall. 1984,32, 1079–1088. [CrossRef]
8.
Arzt, E.; Rösler, J. The kinetics of dislocation climb over hard particles—II. Effects of an attractive particle-dislocation interaction.
Acta Metall. 1988,36, 1053–1060. [CrossRef]
9.
Chinnappan, R. Thermodynamic Stability of Oxide Phases of Fe-Cr Based ODS Steels via Quantum Mechanical Calculations.
Procedia Eng. 2014,86, 788–798. [CrossRef]
10.
Kim, J.H.; Byun, T.S.; Lee, J.H.; Min, J.Y.; Kim, S.W.; Park, C.H.; Lee, B.H. Effects of processing condition on the microstructural
and tensile properties of 14Cr-based oxide dispersion strengthened alloys. J. Nucl. Mater. 2014,449, 300–307. [CrossRef]
11.
Kishimoto, H.; Alinger, M.J.; Odette, G.R.; Yamamoto, T. TEM examination of microstructural evolution during processing of
14CrYWTi nanostructured ferritic alloys. J. Nucl. Mater. 2004,329–333, 369–371. [CrossRef]
12.
Odette, G.R. Recent Progress in Developing and Qualifying Nanostructured Ferritic Alloys for Advanced Fission and Fusion
Applications. Jom 2014,66, 2427–2441. [CrossRef]
13.
Odette, G.R.; Alinger, M.J.; Wirth, B.D. Recent developments in irradiation-resistant steels. Annu. Rev. Mater. Res.
2008
,38,
471–503. [CrossRef]
14.
Tan, L.; He, G.; Liu, F.; Li, Y.; Jiang, L. Effects of Temperature and Pressure of Hot Isostatic Pressing on the Grain Structure of
Powder Metallurgy Superalloy. Materials 2018,11, 328. [CrossRef] [PubMed]
15.
Liu, F.; Wang, Z.; Wang, Z.; Zhong, J.; Zhao, L.; Jiang, L.; Zhou, R.; Liu, Y.; Huang, L.; Tan, L.; et al. High-Throughput
Method–Accelerated Design of Ni-Based Superalloys. Adv. Funct. Mater. 2022. [CrossRef]
16.
Tan, L.; Li, Y.; He, G.; Liu, F.; Nie, Y.; Jiang, L. Optimized hot workability of a powder metallurgy nickel-base superalloy. Mater.
Charact. 2019,147, 340–352. [CrossRef]
17.
Somani, M.C.; Muraleedharan, K.; Prasad, Y.V.R.K.; Singh, V. Mechanical processing and microstructural control in hot working
of hot isostatically pressed P/M IN-100 superalloy. Mater. Sci. Eng. A 1998,245, 88–99. [CrossRef]
18.
He, G.; Liu, F.; Huang, L.; Jiang, L. Hot deformation behaviors of a new hot isostatically pressed nickel based powder metallurgy
superalloy. J. Mater. Res. 2016,31, 3567–3579. [CrossRef]
19.
He, G.; Liu, F.; Huang, L.; Jiang, L. Analysis of forging cracks during hot compression of powder metallurgy nickel-based
superalloy on simulation and experiment. Adv. Eng. Mater. 2016,18, 1823–1832. [CrossRef]
20.
Tan, L.; Li, Y.; Liu, C.; Yang, C.; Ding, H.; Huang, L.; Liu, F.; Qin, Z.; Jiang, L. The evolution history of superalloy powders during
hot consolidation and plastic deformation. Mater. Charact. 2018,140, 30–38. [CrossRef]
21.
Tang, Q.X.; Ukai, S.; Oono, N.; Hayashi, S.; Leng, B.; Sugino, Y.; Han, W.; Okuda, T. Oxide Particle Refinement in 4.5 mass%Al
Ni-Based ODS Superalloys. Mater. Trans. 2012,53, 645–651. [CrossRef]
22.
Lee, M.K.; Park, J.J.; Rhee, C.K. Synthesis and structural properties of Ni-20Cr-Y2O3 nanocomposite alloy prepared by a very
high energy mechanical milling. Mater. Chem. Phys. 2012,137, 129–134. [CrossRef]
Materials 2022,15, 4087 25 of 25
23.
McCarley, J.; Tin, S. Utilization of hot deformation to trigger strain induced boundary migration (SIBM) in Ni-base superalloys.
Mater. Sci. Eng. A 2018,720, 189–202. [CrossRef]
24.
Kumar, S.S.S.; Raghu, T.; Bhattacharjee, P.P.; Appa Rao, G.; Borah, U. Evolution of microstructure and microtexture during hot
deformation in an advanced P/M nickel base superalloy. Mater. Charact. 2018,146, 217–236. [CrossRef]
25.
Konkova, T.; Rahimi, S.; Mironov, S.; Baker, T.N. Effect of strain level on the evolution of microstructure in a recently developed
AD730 nickel based superalloy during hot forging. Mater. Charact. 2018,139, 437–445. [CrossRef]
26.
Huang, K.; Marthinsen, K.; Zhao, Q.; Logé, R.E. The double-edge effect of second-phase particles on the recrystallization
behaviour and associated mechanical properties of metallic materials. Prog. Mater. Sci. 2018,92, 284–359. [CrossRef]
27.
He, G.; Tan, L.; Liu, F.; Huang, L.; Huang, Z.; Jiang, L. Revealing the role of strain rate during multi-pass compression in an
advanced polycrystalline nickel base superalloy. Mater. Charact. 2017,128, 123–133. [CrossRef]
28. Zener, C. Theory of growth of spherical precipitates from solid solution. J. Appl. Phys. 1949,20, 950–953. [CrossRef]
29.
De Castro, V.; Marquis, E.A.; Lozano-Perez, S.; Pareja, R.; Jenkins, M.L. Stability of nanoscale secondary phases in an oxide
dispersion strengthened Fe–12Cr alloy. Acta Mater. 2011,59, 3927–3936. [CrossRef]
30.
Huang, K.; Logé, R.E. A review of dynamic recrystallization phenomena in metallic materials. Mater. Des.
2016
,111, 548–574.
[CrossRef]
31.
Sakasegawa, H.; Ohtsuka, S.; Ukai, S.; Tanigawa, H.; Fujiwara, M.; Ogiwara, H.; Kohyama, A. Particle size effects in mechanically
alloyed 9Cr ODS steel powder. J. Nucl. Mater. 2007,367–370, 185–190. [CrossRef]
32.
Ohtsuka, S.; Ukai, S.; Sakasegawa, H.; Fujiwara, M.; Kaito, T.; Narita, T. Nano-mesoscopic structural characterization of 9Cr-ODS
martensitic steel for improving creep strength. J. Nucl. Mater. 2007,367–370, 160–165. [CrossRef]
33.
Read, H.G.; Murakami, H. Microstructural influences on the decomposition of an Al-containing ferritic stainless steel. Appl. Surf.
Sci. 1996,94–95, 334–342. [CrossRef]
34.
Ukai, S.; Ohtsuka, S. Low cycle fatigue properties of ODS ferritic–martensitic steels at high temperature. J. Nucl. Mater.
2007
,
367–370, 234–238. [CrossRef]
35.
Lee, J.S.; Kimura, A.; Ukai, S.; Fujiwara, M. Effects of hydrogen on the mechanical properties of oxide dispersion strengthening
steels. J. Nucl. Mater. 2004,329–333, 1122–1126. [CrossRef]
36.
Yoshitake, T.; Ohmori, T.; Miyakawa, S. Burst properties of irradiated oxide dispersion strengthened ferritic steel claddings. J.
Nucl. Mater. 2002,307–311, 788–792. [CrossRef]
37.
Ohtsuka, S.; Ukai, S.; Fujiwara, M.; Kaito, T.; Narita, T. Improvement of 9Cr-ODS martensitic steel properties by controlling excess
oxygen and titanium contents. J. Nucl. Mater. 2004,329–333, 372–376. [CrossRef]
38.
Miller, M.K.; Kenik, E.A.; Russell, K.F.; Heatherly, L.; Hoelzer, D.T.; Maziasz, P.J. Atom probe tomography of nanoscale particles
in ODS ferritic alloys. Mater. Sci. Eng. A 2003,353, 140–145. [CrossRef]
39.
Tan, L.; Huang, Z.; Liu, F.; He, G.; Wang, X.; Huang, L.; Zhang, Y.; Jiang, L. Effects of strain amount and strain rate on grain
structure of a novel high Co nickel-based polycrystalline superalloy. Mater. Des. 2017,131, 60–68. [CrossRef]
40.
Wusatowska-Sarnek, A.M.; Miura, H.; Sakai, T. Nucleation and microtexture development under dynamic recrystallization of
copper. Mater. Sci. Eng. A 2002,323, 177–186. [CrossRef]
41.
Xu, H.; Lu, Z.; Ukai, S.; Oono, N.; Liu, C. Effects of annealing temperature on nanoscale particles in oxide dispersion strengthened
Fe-15Cr alloy powders with Ti and Zr additions. J. Alloy. Compd. 2017,693, 177–187. [CrossRef]
42.
Yan, P.; Yu, L.; Liu, Y.; Liu, C.; Li, H.; Wu, J. Effects of Hf addition on the thermal stability of 16Cr-ODS steels at elevated aging
temperatures. J. Alloy. Compd. 2018,739, 368–379. [CrossRef]
43.
Jorge-Badiola, D.; Iza-Mendia, A.; Gutiérrez, I. Study by EBSD of the development of the substructure in a hot deformed 304
stainless steel. Mater. Sci. Eng. A 2005,394, 445–454. [CrossRef]
44.
Bui, Q.H.; Dirras, G.; Ramtani, S.; Gubicza, J. On the strengthening behavior of ultrafine-grained nickel processed from
nanopowders. Mater. Sci. Eng. A 2010,527, 3227–3235. [CrossRef]
45.
Rosenberg, J.; Piehler, H. Calculation of the Taylor Factor and Lattice Rotations for Bcc Metals Deforming by Pencil Glide. Metall.
Trans. 1971,2, 257–259. [CrossRef]
46.
Wasilkowska, A.; Bartsch, M.; Messerschmidt, U.; Herzog, R.; Czyrska-Filemonowicz, A. Creep mechanisms of ferritic oxide
dispersion strengthened alloys. J. Mater. Process. Technol. 2003,133, 218–224. [CrossRef]
47.
Xu, H.; Li, W.; Sha, X.; Meng, J.; Kang, C.; Wang, W.; Zang, X.; Wang, Z. Effects of Zr addition on the microstructural stability of
15Cr-ODS steels under elevated-temperature annealing. Fusion Eng. Des. 2019,138, 231–238. [CrossRef]
48.
Zhang, Y.; Ouyang, T.; Liu, D.; Du, J.; Wang, Y.; Zhang, C.; Suo, J. 9Cr-ODS steel composite material reinforced by Ta layers. J.
Alloy. Compd. 2016,682, 294–301. [CrossRef]
