RETROGRESSION AND RE-AGING OF ALUMINUM ALLOYS (AA 7075) CONTAINING NICKEL
HAIDER T. NAEEMa,b*, KAHTAN S. MOHAMMEDa
aSchool of Materials Engineering, University Malaysia Perils, Taman Muhibbah, 02600 Jejawi, Perlis, Malaysia
bAl-Muthana University, Samawa, Al-Muthana, Iraq
In this study the effects of nickel additions in improving the mechanical properties and microstructure for high-strength aluminum alloys (AA7075) produced by semi-direct chill casting were investigated. Aluminum alloys were homogenized at different temperatures, aged at 120 °C for 24 h (T6), and retrogressed at 180 °C for 30 min and then re-aged at 120 °C for 24 h (RRA). The results of the microstructural analysis showed that adding nickel to aluminum alloy led to form nickel-rich dispersoid particles, such as Al7Cu4Ni, Al4Ni3, Al75Ni10Fe15, Al3Ni2, and Al50Mg48Ni7. These provided particles strengthening of dispersion and fine-grain that led to prevent the recrystallization besides restricted the grain growth. The mechanical properties of the alloys were improved by the strengthened dispersoid particles and precipitation of the matrix base alloy. The highest ultimate tensile strength and Vickers hardness of aluminum alloy containing nickel after the retrogression and re-aging treatment were about 400 MPa and 225 HV, respectively. Microstructure characterization of the alloys were carried out using optical microscopy (OM), scanning electron microscopy (SEM), energy dispersive X-ray (EDX) and X-ray diffraction (XRD).
(Received July 6, 2013; Accepted November 11, 2013)
Keywords: Aluminum alloys (7xxx); Trace element; Re-aging; Microstructural;
Orowan strengthening
1. Introduction
Al-Zn-Mg-Cu aluminum alloys have been studied by many researchers because of their suitable properties and tremendous applications, especially in the aviation and aerospace industries [1].
High-strength aluminum alloys (7xxx) are popular because of their excellent properties, such as high tensile strength, excellent formability, and satisfactory corrosion resistance through heat treatments. The retrogression and re-aging (RRA) is considered one of the most important the heat treatments to produce materials with higher mechanical strength and stress corrosion than aging at T6 temper [2, 3]. The microstructural properties of aluminum–zinc–magnesium–copper alloys after the RRA treatment have been investigated. Typically, the eutectic structures of alloys consist of α-Al and MgZn2. Several other phases such as S-Al2CuMg, T-Al2Mg3Zn3, and Al2Cu are formed by solidifying aluminum alloys (Al-Zn-Mg-Cu) [4, 5]. The zinc–magnesium ratio significantly affects the formation of the MgZn2 phase. High of zinc–magnesium ratio increases the formation other compounds of MgZn2 that strengthen the alloys after heat treatments. Different nucleation agents, such as zirconium (Zr), titanium (Ti), boron (B), and scandium (Sc) have been used in aluminum alloys (7xxx); Ti coupled with B provides the finest microstructure for alloys [6,7]. Various studies have investigated the effect of nickel on the microstructure and mechanical properties of aluminum alloys. Compton et al. [8] found that adding nickel into pure aluminum forms Al3Ni through eutectic reaction and increases alloy hardness. This intermetallic phase is also common in Al-Si-Ni alloys. Yang and Boyuk [9, 10] discovered that blending nickel into
*Corresponding author: [email protected] .
aluminum–silicon alloys forms Al3Ni in addition to Al7Cu4Ni and Al9FeNi during solidification.
For aluminum–silicon–nickel alloys, a nickel-rich dispersion phase insoluble at high temperature improves mechanical properties for applications in high temperatures [11–13]. Moreover, influence addition of nickel on the mechanical properties of Al-Zn-Mg-Cu alloys was investigated using numerous techniques such as; Ingot metallurgy (IM) and rapid solidification (RS).
Shen et al. and Wu et al. [14, 15] found that more semi-coherent and incoherent η′ and η phases in the microstructure of aluminum alloys Al-Zn-Mg-Cu-Ni produced by rapid solidification it's increased hardness after the aging. So a detailed understanding of the microstructural characteristics of Al-Zn-Mg-Cu containing Ni produced by semi-direct chill casting is still lacking.
The purpose of this study aims to determine the effects of nickel additives on microstructural evolutions and mechanical properties of AA7075 aluminum alloys (produced by semi-direct chill casting) after aging and RRA.
2. Experimental procedures 2.1. Research Material
The present study was carried out on AA7075-O aluminum alloy slabs provided by ALCAN GLOBAL AEROSPACE. The slabs were 13 mm thick and 20 mm wide. Nickel of 99%
purity as additives was provided by Merck KGaA. The nominal compositions of the studied alloys are listed in Table 1. The terms “Base alloy” and “Alloy A” refer to as received alloy and alloy with 0.1 wt. % Ni respectively. The chemical composition analysis was carried out using the arc- spark spectrometer (SPECTROMA).
Table 1: The chemical composition of studied alloys (in wt. %).
No. Si Fe Cu Mg Cr Ni Zn Ti Al
Base alloy 0.066 0.24 1.8 2.87 0.18 - 6.68 0.03 balance
Alloy A 0.06 0.22 1.7 2.75 0.18 0.1 6.68 0.028 balance
Alloys were re-melted in a graphite crucible at 850°C (1123 K) in electrical resistance furnace (with accuracy of +/-5°C). The samples were produced by a semi–direct chilling (DC) casting process proceeded in iron steel mold of (150 l x 30 w x 20 h). The mold was preheated to 250°C prior to casting process. The casting speed about 150 mm/min, water flow rate was about 45 l/min and cooling rate of –280⁰C /sec. The alloys were inverted and remelted three times to ensure complete mixing. After the casting; homogenizing treatments conducted for alloys according to step No.1 in Table 2, followed by quenching in cold water immediately after each step of the homogenizing treatments. Thereafter ageing at T6 temper then the retrogression and re-aging (RRA) process their detailed in Table 2; quenching in cold water come after each treat.
Table 2: The homogenizing and heat treatment steps for alloys studied
No. Type Description of treatment
1 Homogenizing 450°C for 2 h+470°C for 24 h+480°C for ½ h
2 Ageing (T6) 120°C for 24 h
3 Retrogression and reaging (RRA)
120°C for 24 h + 180°C for ½h + 120°C for 24h
optical and po of the order t couple
ASTM time. T Each r of the tensile testing were p temper 10mm/
B557M
sample interde equiax
Se an ave (Figs.
alumin direct efficien grain s acted a (a)
2.2. Micro The micro l microscope olished accor average gra to investigate ed energy dis
2.3 Mecha The Hv m M E92-82, “M
To ensure cl reading was specimens.
e test was ca g machine at prepared acc rature on pl
/min., and a M-02a.
3. Resul Fig. 1(a) a es. These sa endritic netw xed grains co
Fig. 1: O
emi-direct ch erage grain s 1a and 1b).
num alloy. A chill castin ncy of the t size of as-qu as a nucleatio
ostructures C structures we e. The specim rding to AST ain size and
e effects of a spersive X-ra anical Testin microhardnes
Mitutoyo DX leanliness th
an average o The highest arried out at t a ram spee
cording to A late tensile a load of 500
lts
and 1(b) show amples conta work of int ontained colu
Optical micros
hill casting fo ize of about Ying et al.
Average grain ng [17]. The technique in
uenched allo on agent on
Characteriz ere analyzed mens were e TM E3-01. T
its analysis additives on t ay (EDX) and ng
ss measurem X256 series”.
e surfaces o of at least ten
t and the lo ambient tem d of 10mm/
ASTM B557 specimens u 0kN. The ten
w the optical ained incipie termetallic c umnar grains
tructures of a
orms fine eq 45 and 39 [16] observ n size is redu e result of s grain refine oy A was re
the grains, in zations d by the optic extracted from They were et
were carried the microstru d X-ray diffr
ments were c . Indentation of the sample n separate m owest values
mperature on /min., and a
7M-02a. Th using INSTR nsile test spe
l micrograph ent dendrite compounds
between the
as-quenched (a
quiaxed grain
m for the b ved a grain s uced by the n semi-direct ement and e efined by ad nhibiting col
(b)
cal microscop m position o tched with K d out using ucture scann raction analy
carried out o n force was s es were polis measurements
of the ten n plate tensi load of 500 he tensile te RON testing ecimens were
hs of as-quen s of alumin around the em.
a) base alloy a
n structures ase alloy and size of abou nucleation ra chill casting nhancement ding nickel lumnar evolu
py (OM) usin f ½ height o Keller’s reage the linear in ing electron ysis (XRD) w
on the speci set to 30N a shed prior to s taken rando reading wer ile specimen kN. The ten st was carri g machine a e prepared a
ched the bas num-rich soli
primary gr
and (b) alloy A
all over the d alloy A sam ut 121 m fo
te during sol g in this stu
of mechani to the alum ution, which
ng Olympus of the ingot, ent. The calc ntercept meth
microscopy were used.
imens accord and 10 sec. d
o Hv measur omly on the re disregarde ns using INS nsile test spe ied out at a at a ram sp according to
se alloy and a lid solution
rains. Sever
A samples
cross section amples, respe or as-cast AA
lidification in udy confirm ical propertie minum alloy.
h usually star
PMG3 ground culation hod. In (SEM)
ding to dwelled
rement.
surface ed. The STRON ecimens ambient peed of
ASTM
alloy A and an ral fine
ns with ectively A 7050
n semi- med the
es. The Nickel rts from
the mo solidif
homog for the structu heat tre
after t respect the T6 additiv and inc compo of hom formed stages.
(a
(a
old wall and fication [18].
Fig. 2:
Fig. 2(a) a genizing and e base alloy a ure was grad
eatments.
Fi Fig. 3(a) a the RRA pr tively. The g 6 temper and ves, which it
creased the s With addit ounds were f mogenization
d compound .
d forms appro
: Optical micr and 2(b) sho d aging treatm and alloy A dually reduce
ig. 3: Optical and 3(b) show
rocess. The grain size of
d RRA treat
’s had higher solidification tion of nicke formed becau n with heat
of dispersed
opriate sites
rostructures of ow the optic ment at T6 te samples, resp ed, and the re
micrographs w the optica average gr the alloy A s tment becau
r density tha n rate.
l, eutectic re use the solub
treatments f d particles res
for the nucl
of (a) base allo cal microgra emper. The a pectively. Th residual phas
of (a) base al al micrograph rain size of samples (Fig use of the in an the matrix eaction create
bility limits o followed by stricted recry
(b
(b
leation of th
oy and (b) allo aphs of the average grai he volume fr ses became s
lloy and (b) al hs of the bas
these alloy gs. 2 and 3b) nteraction of x base alloy a ed dispersion of the nickel quenching ystallization
b
e first alumi
oy A after T6 t alloy sampl n size was a raction of the small and spa
loy A after RR se alloy and ys was abou was signific f matrix base and thus form n compounds
were extend in cold wate and grain gr
inum phases
temper.
les after pro about 50 and e dendritic n arse because
RA.
alloy A spe ut 53 and 4 cantly reduce e alloy with med nucleatio
s in the alloy ded through a
er [19]. The rowth in subs
during
omotive d 41 m network e of the
ecimens 43 m, ed after h nickel on sites y. These a series e newly sequent
quench indicat lattice stoichi 4d). Th X-ray elemen relation Fan [2 and Cu
(a)
F
The scann hed base allo te the non-e
structure ( iometry simi he labeled re
(EDX) scan nts were sim n of Fe was 20], who foun
u.
Fi
(a)
(c)
)
Precipitati
Fig. 4: (a) SEM
ning electron oy sample. Th equilibrium e
(Fig. 4c). G ilar to the T-(
egion in Fig.
in Fig. 4(b) milar to the T
not observe nd that alum
ig. 5: (a) SEM
j
k
on phases
M and (b) EDX
n micrograph he dark area eutectic solid Gray particle (Al2Mg3Zn3) . 4(a) represe
reveals the c T-(AlMg4Zn d in this reg minum alloys
M and (b) EDX
DX scan of as-q
h (SEM) in s indicate the d solution b es (encircled ) phase and t ents the matr chemical com n11) and S-(A gion. These r
(7xxx) conta
X scan of base
(b
(c
quenched bas
Fig. 4 show e primary so etween grain d region (j) the S-(Al2Cu
rix of the ba mposition of Al2CuMg) ph esults are co ain the T, S,
alloy sample
(b)
E(d)
EE
M
b)
)
e alloy sample
ws the micros lid solution a ns; this eute )) were pre uMg) phase w ase alloy. The the labeled r hases. Notab onsistent with
and phase
after T6 temp Element M
Mg K Al K 6 Cu K 2 Zn K
Element Ma Mg K 7.
Al K 61 Si K 0.
Fe K 0.
Cu K 26 Zn K 3.
Element Ma Mg K 5.
Al K 75 Fe K 7.
Cu K 5.
Zn K 6.
Matrix Co le.
structure of and the brigh ectic solution evalent, sug with Al-Cu-F
e energy-dis region, in wh bly, the quan h those obtai es with Al, Z
per.
Mass% At%
3.55 5.4 61.98 77.
29.05 14.
5.43 2.7
ass% At%
.75 13.22 1.63 71.52
.15 0.09 .13 0.07 6.79 13.33
.50 1.68
ass% At%
.32 7.91 5.60 82.81
.59 4.02 .43 2.52 .06 2.74 orre. ZAF
the as- ht areas n had a ggesting Fe (Fig.
spersive hich the ntitative ined by Zn, Mg,
% 40
06 84 70
at 120 with a chemic that of
sample shape spots e AlCuF
(a
Pre
The base a
°C for 24 h script-like s cal composit f the T-AlMg
Fig. 6:
The SEM e after RRA (encircled re exhibit stoich Fecompound
Fig. 7: X cipitation
alloy sample h (T6). The S
shape (encirc tion of the b g4Zn11,S-Al2
(a) SEM and
and EDX sc A. The SEM egion (i), Fig hiometry clo ds (Fig. 6b).
XRD plots of b
i
was homoge SEM in Fig.
cled region ( bright region CuMg, - M
(b) EDX scan
can in Figs. 6 reveals the g. 6c). The ED ose to that of
base alloy as q
enized under 5(a) reveals (i), Fig. 5c).
n. The stoich Mg2Zn11, -M
n of base alloy
6(a) and 6(b existence of DX scan of t f T-AlMg4Zn
quenched, afte
(b
(c
r different tem s the existen
The EDX sc hiometry of t MgZn2, and A
y sample after
b) show the m f precipitates the labeled re n11,S-Al2CuM
er T6 heat trea
b
)
mperatures c ce of precipi can in Fig. 5
his labeled r AlCuFe phas
r RRA heat tre
microstructur s (bright spo egion indicat Mg, -MgZn
atment, and af Element
Mg K Al K Fe K Cu K Zn K Matrix
conditions an itates (bright 5(b) also reve region was c ses.
eatment.
re of the bas ots) with a z ates that these
n2, -Mg2Zn
fter RRA.
Mass% A 3.81 4 87.8 9 0.13 0 1.43 0 6.83 2 Corr. Z
nd aged t spots) eals the close to
se alloy zag-like e bright n11, and At%
4.42 1.93 0.08 0.63 2.94 ZAF
Fig. 7 shows the different X-ray diffraction (XRD) patterns of the base alloy after quenching, T6 temper , and the RRA treatment. The as-quenched base alloy (c) was primarily composed of -(Al), and the secondary phases were T-AlMg4Zn11, S-Al2CuMg, - MgZn2, - Mg2Zn11, and Al23CuFe4, consistent with [20]. The generally accepted precipitation sequences for 7000 series aluminum alloys are as follows: [21-24]: supersaturated solid solution coherent stable Guinier–Preston (GP) zones semi-coherent intermediate (Mg2Zn11) phase incoherent stable (MgZn2) or T(AlMg4Zn11) phase. Metastable phase was the primary precipitation hardening phase of these alloys. The primary precipitations in the matrix were the GP zones and
(Mg2Zn11) phase after aging at 120 °C for 24 h.
The XRD plots in Fig. 7(b); shows that after T6 temper, the base alloy sample exhibited more significant -MgZn2 and -Mg2Zn11 phases, which dissolved during homogenization and then precipitated during aging at T6. No obvious diffraction peak was observed in the Al23CuFe4
phase.
The different X-ray diffraction (XRD) findings for the sample after the RRA process (Fig.
7a) indicated high-intensity diffraction peaks in the MgZn2 and Mg2Zn11 phases. The primary precipitation phases in the matrix (the fine and dispersive GP zone and phase) underwent the RRA treatment after its step first (120°C for 24 h) aging. During retrogression, the GP zones dissolved into the (Mg2Zn11) or -MgZn2 phases. With prolonged retrogression, the undissolved GP zones transformed into the phase and thus formed numerous GP zones and phases that were dissolved in the early stages of another round of retrogression [25, 26]. Finally, abundant nuclei that promote the re-precipitation of the GP zones and phase in the re-aging step were mounted according to the XRD results. The XRD analysis results were consistent with the EDX results.
Fig. 8: (a) SEM and (b) EDX analysis of as-quenched alloy A sample.
The SEM in Fig. 8 (a) shows the microstructure of the as-quenched alloy A sample. The dark areas denote the primary solid solution as indicated in the matrix of the labeled region; the EDX scan points reveal chemical composition close to the T-Al5Mg11Zn4, S-Al2CuMg, -MgZn2,
-Mg2Zn11, and AlCuFe phases (Fig. 8b). The bright areas denote the non-equilibrium solidification eutectic system between grains (Fig. 8c). Gray particles were observed (encircled region (p), Fig. 8d), revealing stoichiometry close to the T-Al5Mg11Zn4,S-Al2CuMg, -MgZn2, - Mg2Zn11, γ-Al-Cu-Ni, and Al-Ni-Fe phases.
Fig. 9: (a) SEM and (b) EDX analysis of alloy A sample after T6 temper.
Fig. 9(a) shows an SEM of the alloy A sample after T6 temper. The encircled region shows a bacillary shape (Fig. 9c). Figure 9(b) reveals chemical composition close to the T- Al5Mg11Zn4,S-Al2CuMg, Al7Cu4Ni, Al50Mg48Ni7, Al4Ni3, Al3Ni2, MgZn2, and Mg2Zn11 phases.
Figure 9(a) shows the prevalence of Ni-rich dispersion particles.
Fig. 10: (a) SEM and (b) EDX analysis of alloy A sample after RRA.
Figure 10 shows the microstructure of the alloy A sample after the RRA process. The bright areas denote newly formed phases in addition to the dispersion particles. Figure 10(a) shows the numerous dispersion particles (see Fig. 10c for highly magnified SEM). Figure 10(b) reveals similar stoichiometry to that of Al4Ni3, Al3Ni2, Al50Mg48Ni7, T-Al5Mg11Zn4, S-Al2CuMg, Al7Cu4Ni, MgZn2, and Mg2Zn11. During RRA, the secondary phase particles were further dissolved into the matrix, and the solubility limit of the nickel additives can be extended with dual solution treatment, consistent with Yuan et al. [25].
Fig. 11: XRD plots for alloy A after quenching, T6 heat treatment, and RRA.
The XRD analysis results for the alloy A samples are shown in Figs. 11(a)–11(c) (RRA, T6, and as-quenched, respectively). The patterns of the as-quenched alloy A sample confirm that the primary eutectic system mainly consisted of (Al), solid solution, and intermetallic compounds (i.e., T-Al5Mg11Zn4, S-Al2CuMg, Al7Cu4Ni, Al50Mg48Ni7, Mg2Zn, Mg2Zn11, Al4Ni3, and Al3Ni2).
The solubility limits of nickel can be extended to form a supersaturated solid solution with the aluminum matrix produced by chill casting and thus form nickel dispersion particles within the aluminum alloy [26]. This result was confirmed by the SEM/EDX and XRD results. Figure 11(b) shows the XRD plots for the alloy A sample after T6 temper and indicates the existence of the Al75Ni10Fe15 phase in addition to the phases in the as-quenched sample (Al50Mg48Ni7, Mg2Zn, Mg2Zn11, Al4Ni3, and Al3Ni2). These dispersive phases had high peaks because of the intensive dissolution of the alloying elements with the nickel additives produced by the homogenization and subsequent heat treatments.
On the other hand, Li et al. [27] found that adding nickel to aluminum alloys Al-Zn-Mg suppresses the formation of the MgZn2 phasein the matrix. These findings contradict the present results according to EDX and XRD analysis (i.e., numerous MgZn2 phases).
Figure 11(a) shows the XRD plots for the alloy A sample after the RRA treatments. The intensity of the diffraction peaks of phases Al4Ni3, Al75Ni10Fe15, Al3Ni2, Al50Mg48Ni7, MgZn2, and Mg2Zn11 increased this because the effects of steps the retrogression and reaging as is detailed above (Fig.10).
alloys 172 an quench of inte
hardne
base a
Fig. 12
Fig. 12 ind with the T6 nd 202 MPa hed alloy sam erdendritic ne
Maximum ess were obta
4. Discu The results alloy and all As-qu T6 RRA
Str ength /MPa
As-q T6 RRA
Hardness (HV)
: Variations in
dicates the s temper and R
(RRA) in u mples after T etworks by th
Fig. 13 ind a
gains of 100 ained.
ussion s indicated t loy A subst uenched A
quenched A
n tensile stren heat t
strength of tw RRA proces ultimate tensi
T6 temper an he heat treatm
dicates the var alloy specimen
0 and 110 HV
hat the yield tantially imp
YS Base 160 250 290
Bas
ngth of alloy sp treatment cond
two alloys u s yielded ma ile strength ( nd RRA was ments (Fig.1
riations in the ns under differ
V (base alloy
d strength (Y proved after
UTS e alloy
180 340 380
se alloy 100 190 210
pecimens befo ditions.
nder differen aximum gain
(UTS). The s due to grai
).
e Vickers hard rent condition
y) and 100 an
YS), UTS, an T6 temper
Y 17 26 30
ore and after d
nt conditions ns of 160 and
variation in in refinemen
dness of the ns.
nd 115 HV (
nd Vickers h and RRA ( YS
AlloyA 72
67 05
Alloy A 110 210 225
different
s. Treating t d 200 MPa (T
the UTS of nt and the ev
(alloy A) in V
hardness of b (in addition
UTS A
198 370 400
the two T6) and the as- volution
Vickers
both the to the
addition of nickel additives in alloy A). The strengthening mechanism of aging at T6 (120 °C for 24 h) for the base alloy is attributed to precipitation hardening; that is, the effects of the GP zones were consistent with the nano-sized metastable precipitates. These precipitates act as pinning points that impede dislocation [23–26]. The XRD and EDX results revealed the existence of and
phases. The YS and UTS of the base alloy sample after RRA significantly improved compared with those of the T6 temper specimens. This improvement is attributed to the partial dissolution of the pre-existing GP zones and phase. The GP zones can act as nucleation sites for particles, and the remaining phase constantly grew during RRA. After the RRA, the solute atoms dissolved in the matrix precipitated again and produced smaller GP zones and phase. Therefore,
(MgZn2) and (Mg2Zn11) phases were more significantly augmented by RRA than by the T6 process. The strengthening mechanisms for nickel microalloying additions to Al-Zn- Mg-Cu alloys can primarily be classified into precipitation strengthening by the alloying element for the base alloy and dispersion and fine-grain strengthening by Ni-rich dispersoid particles. Precipitation was strengthened by the heat treatment. Dispersion can be described as dislocations inhibited by Ni dispersoid particles in the slipping planes. The dispersoid phase particles were looped, bypassed, and/or sheared by dislocation through the Orowan mechanism. The stress required to move a dislocation around a particle is YS, which is increased by dispersion strengthening. Aside from the fine-grain strengthening, Ni-rich dispersoid particles restrict recrystallization and inhibit grain growth (Figs. 2 and 3b). This outcome increases YS.
5. Conclusions
The microstructure of the base alloy contained Al5Mg11Zn4, Al2CuMg, Mg2Zn11, and MgZn2 phases. Adding nickel into the base alloy formed new dispersion particles, such as Al7Cu4Ni, Al4Ni3, Al75Ni10Fe15, Al3Ni2, and Al50Mg48Ni7 particles.
RRA improved the YS of the base alloy more significantly than the T6 temper and significantly increased UTS (i.e., 380 MPa; Vickers hardness = 210).
RRA increased the YS of alloy A more significantly than it did the YS of the base alloy.
The YS of alloy A was increased by the dispersion phase particles (Al7Cu4Ni, Al4Ni3, Al75Ni10Fe15, Al3Ni2, and Al50Mg48Ni7), which restricted recrystallization and grain growth.
The strengthening mechanisms for the two alloys were precipitation related to the alloying elements for the base alloy and the dispersion addition to the fine grains by the Ni-rich dispersoid particles.
Acknowledgements
This work is supported under the grant No. 9001-00338 of the Universiti Malaysia Perlis (UniMAP). The authors gratefully acknowledge the outstanding support provided by the technicians of the work shop of Materials Engineering School, UniMAP.
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