Contents lists available at ScienceDirect
Acta
Materialia
journal homepage: www.elsevier.com/locate/actamat
Full
length
article
Improving
the
high-temperature
oxidation
resistance
of
TiB
2
thin
films
by
alloying
with
Al
Babak
Bakhit
a, ∗,
Justinas
Palisaitis
a,
Jimmy
Thörnberg
a,
Johanna
Rosen
a,
Per
O.
˚A.
Persson
a,
Lars
Hultman
a,
Ivan
Petrov
a, b, c,
J.E.
Greene
a, b, c,
Grzegorz
Greczynski
aa Thin Film Physics Division, Department of Physics (IFM), Linköping University, Linköping SE-58183, Sweden
b Materials Research Laboratory and Department of Materials Science, University of Illinois, Urbana, IL 61801, United States c Department of Materials Science and Engineering, National Taiwan University of Science and Technology, Taipei 10607, Taiwan
a
r
t
i
c
l
e
i
n
f
o
Article history: Received 3 April 2020 Revised 3 July 2020 Accepted 7 July 2020 Available online 12 July 2020 Keywords:
Thin films
Titanium diboride (TiB 2 ) Nanostructure XPS
High temperature oxidation
a
b
s
t
r
a
c
t
Refractorytransition-metal diborides (TMB2) arecandidates forextreme environmentsdueto melting
pointsabove3000°C,excellenthardness,goodchemicalstability,andthermalandelectrical conductiv-ity.However,theytypically sufferfromrapidhigh-temperature oxidation.Here,westudythe effectof Aladditionontheoxidationpropertiesofsputter-depositedTiB2-richTi1-xAlxBy thinfilmsand
demon-stratethat alloying thefilmswith Alsignificantlyincreasesthe oxidationresistancewith aslight de-creaseinhardness.TiB2.4 layers aredepositedbydcmagnetronsputtering (DCMS)fromaTiB2 target,
whileTi1-xAlxByalloyfilmsaregrownbyhybridhigh-powerimpulseanddcmagnetronco-sputtering
(Al-HiPIMS/TiB2-DCMS).Allas-depositedfilmsexhibitcolumnarstructure.Thecolumnboundaries ofTiB2.4
areB-rich, while Ti0.68Al0.32B1.35 alloys have Ti-rich columns surroundedby a Ti1-xAlxBy tissue phase
whichispredominantlyAlrich.Air-annealingTiB2.4 attemperaturesabove500 °Cleadstothe
forma-tionofoxidescales thatdo not containB and mostlyconsistofarutile-TiO2 (s)phase.Theresulting
oxidationproductsarehighlyporousduetotheevaporationofB2O3(g)phaseaswellasthecoarsening
ofTiO2 crystallites.Thispooroxidationresistanceissignificantlyimprovedbyalloying withAl.While
air-annealingat800°Cfor0.5hresultsintheformationofan~1900-nmoxidescaleonTiB2.4,the
thick-nessofthescaleformedontheTi0.68Al0.32B1.35 alloysis~470nm.Theenhancedoxidationresistanceis
attributedtotheformation ofadense,protectiveAl-containingoxidescalethatconsiderablydecreases theoxygendiffusionratebysuppressingtheoxide-crystallitescoarsening.
© 2020ActaMaterialiaInc.PublishedbyElsevierLtd. ThisisanopenaccessarticleundertheCCBYlicense.(http://creativecommons.org/licenses/by/4.0/)
1. Introduction
Refractory transition-metal diborides (TMB 2), classified as ultra- high temperature ceramics, are promising materials for extreme thermal and chemical environments which include, individually or in combination, temperatures above 20 0 0 °C, drastic chemi- cal reactivity, hydrostatic pressure, mechanical stress, wear, and very high levels of radiation and heat gradients [1, 2]. Their high strength at elevated temperatures together with thermal conduc- tivity, which provides a high thermal-shock resistance in severe heat fluxes, make TMB 2 ceramics suitable for aerospace applica- tions such as rocket components, atmospheric reentries, jet en- gine turbines, propulsion systems, and sharp leading edges in hy- personic vehicles, with speeds exceeding Mach 5 [1–13]. In addi- tion, there is also a growing demand for employing TMB 2in high-
∗ Corresponding author.
E-mail address: babak.bakhit@liu.se (B. Bakhit).
temperature electrodes [1, 14, 15], molten metal containment [1], refractory crucibles [16], thermocouple protection tubes in steel baths and aluminum reduction cells [17, 18], reinforcement fibers [19, 20], solar power [21, 22], and armor applications [18]. The use of diborides in advanced nuclear fission reactors as neutron ab- sorber materials, which are intentionally introduced into reactor cores to control the neutron balance, is also of great interest [17, 23].
TMB 2 have recently received increasing attention as the new class of hard ceramic protective thin films. They are already being coated on cutting tools [24–29], engine components in aerospace applications [11, 30], nuclear fusion devices [31, 32], as well as metallic contacts in semiconductor [33] and microelectronic com- ponents [34, 35]. These films typically crystallize in a hexagonal AlB 2 crystal structure (P6/mmm, SG-191) in which B atoms form graphite-like honeycomb sheets between hexagonal-close-packed TM layers [36, 37]. The strong covalent bonding between TM and B atoms as well as within the honeycomb B sheets provides high https://doi.org/10.1016/j.actamat.2020.07.025
melting temperature, hardness, and stiffness [38, 39], while metal- lic bonding within TM layers results in good electrical and ther- mal conductivities for these compounds [36]. Hence, this unique combination of ceramic and metallic properties makes TM diboride thin films promising candidates for many applications.
However, sputter-deposited TMB 2 films typically contain excess B, with B/TM ratio ranging from 2.4 to 3.5 [1, 40–47]. TM diborides, contrary to TM nitrides which have very wide single-phase regions [48, 49], are line-compounds [17] for which deviations from sto- ichiometry, B/TM ratio = 2, can lead to the formation of second phases. Thus, controlling the composition of diboride layers is cru- cial as it has significant influences on film properties. Among dif- ferent approaches suggested for obtaining stoichiometric sputter- deposited TMB 2 [43, 44], growing the layers by either dc mag- netron sputtering (DCMS) equipped with Helmholtz coils [45] or high-power impulse magnetron sputtering (HiPIMS) [56, 47]leads to controllably tuning the B/TM ratio. Both techniques exploit dif- ferences in the ionization probabilities between TM and B atoms in order to guide ion fluxes to the substrate. Another critical is- sue which restricts the applications of TMB 2 films is their inher- ent brittleness [50]. We recently proposed a strategy to simultane- ously increase both hardness and toughness of ZrB 2 layers by al- loying with Ta [51]. The hardness of the alloys grown by hybrid Ta-HiPIMS/ZrB 2-DCMS co-sputtering increases from ~35.0 GPa for ZrB 2.4 to ~42.0 GPa for Zr 0.7Ta 0.3B 1.5, accompanied by an enhance- ment in toughness from 4.0 to 5.2 MPa √ m, due to a corresponding transition in nanostructure from B- to Ta-rich column boundaries together with a decrease in column sizes [51].
Moreover, TMB 2 suffer from poor high-temperature oxidation resistance, which is a critical property for many aggressive envi- ronments. In general, the oxidation products of bulk monolithic TMB 2, mostly synthesized by powder metallurgy processes [11, 52, 53], are typically TMO 2 ( s) and glassy B 2O 3 ( l) phases, with lim- ited intermixing [54]. TMB 2start to oxidize at temperatures below ~450 °C [1, 19], depending on their structure, chemical composi- tion, oxygen partial pressure, and oxidation temperature T a [1, 7, 55–57]. For T a ≤ 10 0 0 °C and constant oxygen partial pressure, the oxide scale is composed of a porous, but stable crystalline TMO 2 ( s) skeleton filled by an amorphous B 2O 3 ( l) phase [54], which is highly corrosive [19]. Since the B 2O 3 ( l) phase has a higher wet- tability than TMO 2 ( s) [58, 59], a continuous B 2O 3 ( l) layer can form on the surface of the scale [60, 61]. Oxidation kinetics in this regime are limited by the oxygen diffusion through the B 2O 3 ( l) phase [62, 63]. This continuous layer can behave as an oxidation barrier which maintains its protective effect up to T a = ~10 0 0 °C [54, 55, 64]. However, the oxidation behavior changes at T a > 10 0 0 °C; rather than forming a stable passive oxide scale, B 2O 3 ( l) rapidly evaporates which results in the formation of a porous oxide scale that does not passivate the surface [7, 55, 64].
Many attempts have been made to enhance the oxidation re- sistance of the bulk refractory diborides, mostly by alloying with Cr, Al, and Si or adding secondary compounds such as SiC, TMSi 2 (TM = Zr, Mo, Ta, W), Si 3N 4, and ZrC [1, 7, 10, 56, 58, 64–71]. The high-temperature oxidation rate of bulk TMB 2 which contain ad- ditives is significantly lower than monolithic TMB 2. For example, adding ~20% SiC can effectively improve the oxidation resistance of ZrB 2 due to the formation of a protective oxide scale which is composed of a SiO 2-rich outer layer and a ZrO 2-rich inner layer [56, 64, 72].
The structural and mechanical properties of Ti 1-xAl xB y layers were previously investigated [73–77]; however, little is known about the oxidation resistance of these films. Here, we study the effect of Al addition on the oxidation properties of TiB 2-rich Ti 1-xAl xB ythin films, compared to TiB 2.4films (TiB 2with 0.4 excess B) deposited by DCMS. We use DCMS from a TiB 2 target in pure Ar atmosphere to grow TiB y films at ~475 °C and vary the layer
composition by co-sputtering with partially ionized Al fluxes from the Al-HiPIMS source to obtain Ti 1-xAl xB y alloy films (a hybrid Al-HiPIMS/TiB 2-DCMS technique [79]). The substrate bias is syn- chronized to the Al-rich portion of each HiPIMS pulse [80], based on the input from ion mass spectrometry analyses conducted at the substrate position [79]. Out of all compositions investigated, Ti 0.68Al 0.32B 1.35alloys are chosen for further nanostructural and ox- idation studies as they are characterized by hexagonal structure, relatively low residual stress, high hardness and good indentation toughness. The column boundaries of TiB 2.4 are B-rich, while the Ti 0.68Al 0.32B 1.35alloy films have Ti-rich columns surrounded by an Al-rich Ti 1-xAl xB y tissue phase which is B deficient. Compared to TiB 2.4, the Ti 0.68Al 0.32B 1.35 alloys show significantly better high- temperature oxidation resistance.
2. Experimental
Ti 1-xAl xB yalloy films, in which x=Al/(Ti +Al) and y=B/(Ti +Al), are grown in a CC800/9 CemeCon AG sputtering system [81] equipped with rectangular 8.8 × 50 cm2 stoichiometric TiB
2 and elemental Al targets. Al 2O 3(0 0 01) and Si(001), 1.5 × 1.5 cm 2, substrates are cleaned sequentially in acetone and isopropyl alco- hol, and then mounted symmetrically with respect to the targets, tilted toward the substrates, resulting in a 21 ° angle between the substrate normal and the normal to each target. The Al 2O 3(0 0 01) substrates are used for residual stress and nanoindentation mea- surements, while the Si(001) substrates are used for nanostructural and oxidation studies. The target-to-substrate distance is 20 cm, and the system base pressure is 3.0 × 10 −6 Torr (0.4 mPa). The chamber is degassed before deposition by applying 8.8 kW to each of two resistive heaters for 2 h, resulting in a temperature of ~475 °C at the substrate position. The total Ar pressure during de- position is 3.0 mTorr (0.4 Pa). Prior to the deposition process, the targets are DCMS pre-sputtered in Ar at 2 kW for 60 s with closed cathode shutters.
TiB y films are grown by DCMS at a TiB 2-target power P TiB2 of 40 0 0 W and a negative dc substrate bias of 200 V. For Ti 1-xAl xB y deposition, a hybrid Al-HiPIMS/TiB 2-DCMS scheme is used in which the Al magnetron is operated in HiPIMS mode, with 30-μs pulses and 500-Hz frequency, to supply pulsed energetic Al + fluxes, while the TiB 2 target is continuously sputtered by DCMS. The average power applied to the HiPIMS Al target is maintained constant at 1500 W. The Al-target peak current density during the HiPIMS pulse is ~0.91 A/cm 2. A pulsed negative substrate bias of 200 V is applied in synchronous with the 100-μs Al-ion-rich por- tion of each HiPIMS pulse, starting 30 μs after the target-pulse on- set. At all other times, the substrates are at a negative floating potential of 10 V, in order to reduce the Ar incorporation in the films [79]. While all deposition parameters are maintained con- stant, the power P TiB2 applied to the DCMS TiB 2 target is changed
from 30 0 0 to 50 0 0 W in increments of 50 0 W in order to change the Al/(Ti +Al) ratio, x, in the films.
Surface and fracture cross sections of as-deposited and air- annealed films are examined using a Zeiss LEO 1550 scanning electron microscope (SEM). X-ray diffraction (XRD)
θ
−2θ
scans are carried out using a Philips X ´Pert X-ray diffractometer with a Cu K α source (λ
= 0.15406 nm) in order to determine crys- tal structure and orientation. Cross-sectional and plan-view trans- mission electron microscopy (TEM) analyses are carried out in a double C s aberration-corrected FEI Titan 3 60–300 electron mi- croscope operated at 300 kV; Z-contrast images are obtained in scanning TEM high-angle-annular-dark-field (STEM-HAADF) mode. Energy-dispersive X-ray (EDX) and electron energy-loss spec- troscopy (EELS) elemental maps are also acquired using the Su- perX and GIF Quantum ERS spectrometers embedded in the FEI in-strument. TEM specimens are prepared by focused ion beam (FIB) technique employing a Carl Zeiss Cross-Beam 1540 EsB system.
Substrate wafer curvatures are measured to determine the in- plane residual stress of the as-deposited films on Al 2O 3(0 0 01), based on the modified Stoney equation [82, 83]:
σ
f=Msh2s
/
(
6Rshf)
,where
σ
f is the average biaxial stress; h fand h sare film and sub- strate thicknesses, respectively; R sis the substrate radius of curva- ture; and M s is the substrate biaxial modulus, which is 602 GPa for Al 2O 3 [84]. Substrate curvatures are determined from rocking- curve measurements using a PANalytical Empyrean high-resolution X-ray diffractometer operated at 45 kV and 40 mA. Reportedσ
f values are corrected for thermal stresses due to cooling the sam- ples from T sto room temperature [85].X-ray photoelectron spectroscopy (XPS) is used to analyze the elemental distribution and chemistry of the layers in the as- deposited state as well as after annealing in air. Analyses are performed in a Kratos Axis Ultra DLD instrument employing monochromatic Al K α radiation (h
ν
= 1486.6 eV) and operating with a base pressure lower than 1.1 × 10−9 Torr (1.5 × 10−7 Pa) during spectra acquisition. XPS depth profiles are acquired by sputter-etching with 0.5-keV Ar + ions incident at an angle of 70 ° with respect to the sample surface normal. The low Ar + energy and shallow incidence angle are chosen in order to minimize the effect of sputtering damage on core level spectra [86, 87]. Sample areas analyzed by XPS are 0.3 × 0.7 mm 2and located in the center of 3 × 3 mm 2ion-etched regions. The binding energy scale is cal- ibrated using the ISO-certified procedure [88], and the spectra are referenced to the Fermi edge in order to avoid uncertainties associ- ated with employing the C 1s peak from adventitious carbon [89]. XPS depth scales are converted from time to distance (nm) by us- ing the sputter-etching rate of TiB 2.4 layers (0.68 nm/min) and the average film thickness measured by SEM.Film compositions are determined by time-of-flight elastic re- coil detection analysis (ToF-ERDA) in a tandem accelerator. ToF- ERDA is carried out with a 36 MeV 127I 8+ probe beam incident at 67.5 ° with respect to the sample surface normal and recoils are de- tected at 45 °. Nanoindentation analyses of the layers are performed in an Ultra-Micro Indentation System with a sharp Berkovich dia- mond tip calibrated using a fused-silica standard sample. For hard- ness H measurements, the load is increased from 5 to 25 mN at increments of 0.5 mN, and the results are analyzed using the Oliver and Pharr method [90]. Indents to depths ≥ 10% of the film thickness are excluded in the analysis. Film responses to high local stresses induced by a diamond cube-corner tip, known as nanoin- dentation toughness, are studied by measuring the average lengths of radial cracks around sample indents.
TiB 2.4 and Ti 0.68Al 0.32B 1.35 specimens, 1.0 × 0.5 cm 2, are an- nealed at temperatures T aup to 800 °C in air for times t aranging from 0.5 to 8.0 h using a high-temperature furnace from MTI Cor- poration (GSL-1100 ×-S). The heating rate is constant at 10 °C/min, and the specimens are cooled down to room temperature, while the furnace is turned off.
3. Results
3.1. Compositionandnanostructure
Variations in the x and y values of as-deposited Ti 1-xAl xB y thin films, determined by ToF-ERDA, are plotted in Fig. 1 as a func- tion of P TiB2. TiB y films grown using DCMS ( P TiB2 = 40 0 0 W) are overstoichiometric as the B/(Ti +Al) ratio y= 2.4. The Al/(Ti +Al) ra- tio, x, in the Ti 1-xAl xB yalloys deposited by hybrid Al-HiPIMS/TiB 2- DCMS co-sputtering decreases from 0.35 for P TiB2 = 30 0 0 W, to 0.32 for P TiB2 = 3500 W, to 0.29 for 40 0 0 W ≤ PTiB2 ≤ 500 0 W,
Fig. 1. Variations in the x and y values of as-deposited Ti 1-x Al x B y thin films as a function of TiB 2 target power P TiB2 .
Fig. 2. XRD θ−2 θ scans of as-deposited (a) TiB 2.4 , (b) Ti 0.71 Al 0.29 B 1.54 , (c) Ti 0.71 Al 0.29 B 1.51 , (d) Ti 0.71 Al 0.29 B 1.45 , (e) Ti 0.68 Al 0.32 B 1.35 , and (f) Ti 0.65 Al 0.35 B 1.30 thin films. The peak at 32.8 ° arises from the forbidden 002 Si(001) substrate reflection.
while y gradually increases from 1.30 to 1.35 to 1.45 to 1.51 to 1.54 as a function of P TiB
2. Ar concentration is ~1.2 at.% in the TiB 2.4
films, while it is ~0.5 at.% for all alloys. The detailed elemental compositions of the as-deposited films are given in supplementary Table I.
Fig. 2 shows XRD
θ
−2θ
scans of as-deposited Ti 1-xAl xB y thin films grown on Si(001) substrates. Vertical solid and dashed lines correspond to reference powder-diffraction peak positions for TiB 2 [91]and AlB 2 [92], respectively. The peak at 32.8 ° arises from the forbidden 0 02 Si(0 01) substrate reflection that appears due to mul- tiple scattering events [93]. All reflections in the XRD patterns of the as-deposited Ti 1-xAl xB y films with x≤ 0.32 originate from the hexagonal crystal structure, while the pattern of Ti 0.65Al 0.35B 1.30 grown at P TiB2 = 30 0 0 W has two extra X-ray peaks indicating the presence of an additional intermetallic TiAl phase. The positions of 0 01 and 0 02 reflections shift toward lower 2θ
values with increas- ing x, corresponding to an increase in the out-of-plane c lattice pa- rameter from 0.321 nm for TiB 2.4 to 0.328 nm for Ti 0.65Al 0.35B 1.30. All films exhibit 001 fiber texture in which the 001 and 002 peaks are dominant, with a minor 101 component that becomes stronger as a function of Al content.After compensation for the thermal stresses, residual stresses for all layers grown on Al 2O 3(0 0 01) are compressive with
Fig. 3. Typical cross-sectional and plan-view STEM images, with inset corresponding high-resolution STEM images and SAED patterns, of as-deposited (a and c) TiB 2.4 and (b and d) Ti 0.68 Al 0.32 B 1.35 thin films. The XSTEM SAED patterns are acquired from the areas indicated by circles in (a) and (b). The region indicated by a dashed box in inset (b) exhibits the area where EDX elemental maps, shown in Fig. 5 , are acquired from.
σ
f = ~0.9 GPa for TiB 2.4, ~0.4 GPa for Ti 0.71Al 0.29B 1.54,~0.2 GPa for Ti 0.71Al 0.29B 1.51, and ~0.1 GPa for Ti 0.71Al 0.29B 1.45, Ti 0.68Al 0.32B 1.35, and Ti 0.65Al 0.35B 1.30. The nanoindentation hardness of TiB 2.4 thin films is 42 ±3 GPa, while all alloys have lower hardness values be- tween 35 ±2 GPa and 39 ±3 GPa. Indenting the films by a cube- corner indenter with a load of 100 mN results in the formation of radial cracks and spallation, which is a common feature of brittle materials, in the films with y≥ 1.51. The average length of radial cracks around the cube-corner indents decreases by increasing Al content. The procedure of determining the indentation toughness and typical SEM images from their cube-corner-indented surfaces are presented in supplementary Fig. S1. The as-deposited Ti 1-xAl xB y alloys with y ≤ 1.45 show higher indentation toughness than the other films. Out of all compositions investigated, Ti 0.68Al 0.32B 1.35al- loys are chosen for nanostructural and oxidation studies as they are characterized by hexagonal structure (as revealed by XRD), rel- atively low residual stress, good indentation toughness, and high hardness (39 ±3 GPa).Fig. 3compares cross-sectional and plan-view STEM images of the as-deposited TiB 2.4 and Ti 0.68Al 0.32B 1.35 thin films, with corre- sponding inset selected-area electron diffraction (SAED) patterns. The cross-sectional STEM (XSTEM) micrographs in Fig. 3(a) and 3(b) reveal that the as-deposited films consist of dense columns with no discernable porosity and open boundaries. The thick- ness of as-deposited TiB 2.4 is ~1400 nm, while the as-deposited Ti 0.68Al 0.32B 1.35 films are ~1600-nm thick. Both layers exhibit a competitive growth in which fine columns form near the sub- strate, while fewer columns which become wider extend along the growth direction. The SAED patterns obtained from near the substrates, insets in Fig. 3(a) and 3(b), are composed of weak
diffraction arcs with 001, 101, and 002 components along the growth direction, consistent with the XRD results in Fig.2. As the film thickness increases, the 001 oriented columns begin to be- come dominant in a competitive columnar growth. The plan-view SAED patters acquired from the top part of the films, insets in Fig. 3(c) and 3(d), completely lack the 001 ring which is an in- dication that this set of 001 planes are normal to the electron beam direction, revealing that the films have a strong 001 tex- ture. The TiB 2.4 columns are slightly inclined with respect to the substrate surface normal, Fig.3(a), due to the 21 ° deposition an- gle between the substrate and the TiB 2 target. The top surface of the alloys exhibits faceted columns and a higher surface rough- ness compared to TiB 2.4, in agreement with their SEM surface morphologies shown in supplementary Fig. S2. In addition, higher resolution XSTEM images, insets in Fig. 3(a) and 3(b), show that the columns of Ti 0.68Al 0.32B 1.35 are wider than TiB 2.4. The column boundaries of as-deposited Ti 0.68Al 0.32B 1.35 appear dark indicating a lower average mass than that of the adjacent columns, inset in Fig.3(b).
The Z-contrast plan-view image of as-deposited TiB 2.4, Fig.3(c), shows a nanostructure with no porosity and an average column width of ~7 nm. There is an asymmetry in the columns shape; the columns are elongated toward the incoming flux from the TiB 2tar- get, consistent with the XSTEM results in Fig.3(a). The higher reso- lution image shown as inset in Fig.3(c) reveals that the films con- sist of nanocolumns with a contrast difference between brighter columns and darker column boundaries. The dark regions corre- spond to low-Z column-boundary areas as observed in the over- stoichiometric diboride layers [51, 74, 94, 95]. The Z-contrast plan- view image of as-deposited Ti 0.68Al 0.32B 1.35 in Fig. 3(d) exhibits
Fig. 4. Plan-view (a) STEM Z-contrast image and (b) Ti EDX map with (c) corresponding EELS spectra from the columns and column boundaries of as-deposited TiB 2.4 grown by DCMS. Plan-view (d) STEM Z-contrast image, and (e) Ti and Al EDX maps with (f) corresponding EELS spectra from the columns and column boundaries of as-deposited Ti 0.68 Al 0.32 B 1.35 grown by hybrid Al-HiPIMS/TiB 2 -DCMS co-sputtering.
wide columns, ~40 nm, with dark regions appeared both inside the columns and in the column boundaries.
Fig. 4 is comprised of STEM Z-contrast plan-view images, with corresponding EDX elemental maps and EELS spectra, of as- deposited TiB 2.4 and Ti 0.68Al 0.32B 1.35thin films. The Ti EDX map in Fig.4(b) shows that the dark column-boundary areas in the TiB 2.4 Z-contrast image, Fig.4(a), are Ti deficient, while the correspond- ing EELS spectra in Fig.4(c) confirm that the column boundaries are B rich. However, the Z-contrast image of Ti 0.68Al 0.32B 1.35reveals dark regions both in the columns and column boundaries, Fig.4(d). The Ti and Al EDX maps, shown in Fig.4(e), affirm that there are local changes in Ti and Al concentrations, as previously reported by Nedfors et al. [77] for sputter-deposited (Ti,Al)B 2 alloys. The EELS spectra in Fig.4(f) show that the Ti 0.68Al 0.32B 1.35 column bound- aries are B deficient with respect to the columns. In general, the column boundaries of TiB 2.4films grown by DCMS are B-rich, while the Ti 0.68Al 0.32B 1.35 alloys have Ti-rich columns surrounded by an Al-rich Ti 1-xAl xB ytissue phase which is B deficient.
Fig. 5 shows the XSTEM image, with corresponding Ti and Al EDX maps, of as-deposited Ti 0.68Al 0.32B 1.35 acquired from the re- gion indicated by a dashed box in the inset of Fig.3(b). The EDX maps reveal that the amount of Al in the column boundaries is sig- nificantly higher than in the columns. Compared to the Ti 1-xAl xB y alloy films sputter-deposited by DCMS [74–77], the pronounced Al segregation in Ti 0.68Al 0.32B 1.35 column boundaries can be at- tributed to the enhanced adatom mobility caused by the Al + ion bombardment. We previously reported a related nanostructure for the Zr 1-xTa xB y thin films, with x ≥ 0.2, deposited by hybrid Ta- HiPIMS/ZrB 2-DCMS co-sputtering [51]. These films showed a self- organized columnar core/shell nanostructure in which crystalline hexagonal Zr-rich stoichiometric Zr 1-xTa xB 2 cores are surrounded by narrow dense, disordered Ta-rich shells that are B deficient. The disordered shells have the structural characteristics of metallic- glass thin films which exhibit both high strength and toughness. Hence, such a nanostructure combines the benefits of crystalline diboride columns, providing the high hardness, with the dense
Fig. 5. (a) High-resolution XSTEM image, (b) with corresponding Ti and Al EDX maps, of as-deposited Ti 0.68 Al 0.32 B 1.35 acquired from the region indicated by a dashed box in the inset of Fig. 3 (b).
metallic-glass-like shells which give rise to enhanced toughness [78,51].
3.2.Annealinginair
The TiB 2.4 thin films are air-annealed at different temperatures T a for t a = 1.0 h, and the average B concentrations in their oxida- tion products are determined from XPS depth-profile data, shown in Fig.6. The XPS data are normalized to the ToF-ERDA composi- tions. The B concentration in the oxide scale of TiB 2.4 thin films
Fig. 6. XPS-obtained B concentrations in the oxidation products of the TiB 2.4 thin films air-annealed at different temperatures T a for t a = 1.0 h. The concentrations are normalized to the compositions obtained from ToF-ERDA.
Fig. 7. XRD θ−2 θscans of (a) TiB 2.4 and (b) Ti 0.68 Al 0.32 B 1.35 thin films air-annealed at T a = 700 °C for t a = 1.0 h. The peak at 32.8 ° arises from the forbidden 002 Si(001) substrate reflection.
decreases from ~22 at.% for T a = 300 °C to ~5 at.% for T a = 500 °C. The oxidation products of the films air-annealed at T a> 500 °C are highly B deficient (~0 at.%).
To investigate the effect of Al addition on the high-temperature oxidation resistance, TiB 2.4 and Ti 0.68Al 0.32B 1.35 layers are air- annealed at 700 °C, the temperature at which the oxide scale of TiB 2.4 contains ~0 at.% B. Fig.7compares the XRD
θ
−2θ
scans of TiB 2.4 and Ti 0.68Al 0.32B 1.35 thin films air-annealed at T a = 700 °C for t a = 1.0 h. Some extra reflections appear in the XRD pattern of air-annealed TiB 2.4, Fig. 7(a), which can be assigned to B 2O [96] and rutile-TiO 2 [97] phases. In contrast, there are no addi- tional high-intensity reflections in the air-annealed Ti 0.68Al 0.32B 1.35 XRD pattern, except at 25.2 ° where a low-intensity, broad peak arises, Fig. 7(b). This signal can be attributed to B 2O [96] and Al 8B 2O 15[98]phases. The intensity of the X-ray reflections arising from the hexagonal structure in the as-deposited TiB 2.4 decreases after annealing, while the corresponding peaks in the air-annealed Ti 0.68Al 0.32B 1.35 XRD pattern have significantly higher intensities than in the as-deposited alloys, Fig.2. It indicates an increase inFig. 8. (a) Ti 2p, (b) B 1 s, and (c) O 1 s XPS core-level spectra acquired from the TiB 2.4 thin films air-annealed at T a = 700 °C for t a = 1.0 h as a function of sputter- ing depth d . (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
the crystallinity of Ti 0.68Al 0.32B 1.35 films after 1.0 h annealing at 700 °C. In addition, the FWHM of the 001 reflection from the TiB 2.4 XRD pattern decreases after annealing from 1.1 ° to 1.0 °, while an increase from 0.4 ° to 0.5 ° takes place for Ti 0.68Al 0.32B 1.35. There is also a slight shift in the positions of 00l reflections for both layers toward higher 2
θ
values due to annealing which can be attributed to releasing residual stresses.To evaluate the chemistry of the oxide scales, the primary XPS core-level spectra acquired from the TiB 2.4and Ti 0.68Al 0.32B 1.35thin films air-annealed at T a = 700 °C for t a = 1.0 h are plotted as a function of sputtering depth d in Figs. 8 and9, respectively. For completeness, O 1s spectra are also included. The Ti 2p spectra obtained from d ≥ 22 nm, shown in Fig. 8(a), have broad and convoluted signals consisting of two dominant doublets. These sig- nals arise from Ti bound to O in TiO 2and sub-stoichiometric TiO (
< 2) [99–101]. The formation of the sub-stoichiometric TiO phase is due to the induced damage during the surface sputter-
Fig. 9. (a) Ti 2p, (b) Al 2p, (c) B 1 s, and (d) O 1 s XPS core-level spectra acquired from the Ti 0.68 Al 0.32 B 1.35 thin films air-annealed at T a = 700 °C for t a = 1.0 h as a function of sputtering depth d . (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
etching process [101, 102]. Ti 2p 3/2signals from the TiO 2 and sub- stoichiometric TiO phases appear at 459.0 and 455.3 eV, respec- tively. No B signal is detected for d ≤ 435 nm, Fig. 8(b). Fig.8(c) shows O 1s spectra that consist of convoluted signals; the domi- nant ones appear at 530.8 eV that is characteristic of TiO 2[99, 100, 103–105], while the signals assigned to sub-stoichiometric TiO are shifted to higher BE as the valence charge density on corre- sponding O atoms is lower than on those in TiO 2 [101].
Fig. 9 shows the Ti 2p, Al 2p, B 1s, and O 1s XPS core-level spectra acquired from the Ti 0.68Al 0.32B 1.35 thin films air-annealed at T a = 700 °C for t a = 1.0 h as a function of d. For 11 nm ≤
d < 272 nm, the Ti 2p spectra is essentially identical to that of TiB 2.4 with signals coming from Ti in TiO 2 and sub-stoichiometric TiO , Fig.9(a). The spectra for d≥ 272 nm consist of an additional doublet with Ti 2p 3/2component at 454.4 eV, assigned to Ti in Ti- Al-B. The oxide signals disappear for d> 316 nm, and the Ti-Al-B signals become dominant in the spectra. Sets of Al 2p core-level spectra in Fig.9(b) show broad signals at ~75.6 eV for d≤ 272 nm assigned to Al-O bond [106]. The signal intensity decreases consid- erably for d ≥ 316 nm, consistent with changes observed in cor- responding Ti 2p spectra, and a new peak appears at 74.2 eV due to Al in Ti-Al-B. The B 1s spectra, Fig.9(c), for d< 272 nm reveal broad peaks centered at ~193.5 eV that can be assigned to B in B- O [107]. Starting from d= 272 nm, in addition to the B-O peak, an extra signal appears at 187.7 eV, indicative of B in Ti-Al-B [108]. The intensity of the B-O signal drops significantly for d> 272 nm, while that of the latter peak increases such that the B 1s spec- tra for d> 316 nm consist of a single Ti-Al-B signal, in agreement with the changes in the Ti 2p and Al 2p signals shown in Fig.9(a) and 9(b). The O 1s spectra in Fig.9(d) has convoluted signals aris- ing from TiO 2 (530.0 eV), sub-stoichiometric TiO (531.1 eV), Al-O (533.0 eV), and B-O (532.0 eV) for d ≤ 316 nm [101, 102, 106, 107]. The relative contributions of different oxides exhibit slight varia- tions as a function of depth which are fully consistent with corre- sponding Ti 2p, Al 2p, and B 1s spectra. For example, an increase
in the Al-O and B-O component of the O 1s spectra at d= 142 nm is consistent with an increased intensity of both Al 2p and B 1s peaks (not the case for Ti 2p peak). Furthermore, the drop in Al 2p intensity observed at d= 272 nm is responsible for the decreased intensity of the corresponding component in the O 1s spectra. For
d> 272 nm, the O 1s intensity drops quickly which indicates the oxide/film interface. The overall evolution of XPS core level signals as a function of depth reveals that the Ti 0.68Al 0.32B 1.35 alloys are oxide free for d≥ 300 nm.
XPS depth profiles reconstructed from the raw spectra of the TiB 2.4 and Ti 0.68Al 0.32B 1.35 thin films air-annealed at T a = 700 °C for t a = 1.0 h are compared in Fig.10. XPS values are normalized to the ToF-ERDA compositions in order to minimize the influence of preferential sputtering effects that cannot be completely avoided during depth profiling with Ar +ions [86]. The depth profile of air- annealed TiB 2.4, Fig. 10(a), reveals that the thickness of the oxide scale exceeds 450 nm. Moreover, the oxidation product does not contain B up to d= ~450 nm and consists exclusively of ~65 at.% O and ~35 at.% Ti, indicating TiO 2 formation. In contrast, the av- erage thickness of the oxide scale formed on the Ti 0.68Al 0.32B 1.35 alloys, Fig.10(b), is ~300 nm. Contrary to the oxidized TiB 2.4, the Ti 0.68Al 0.32B 1.35 scale contains ~10 at.% B. Overall, the XPS depth profiles show good agreement with the depth profiles obtained from ToF-ERDA (supplementary Fig. S3).
The cross-sectional and plan-view STEM images of TiB 2.4 and Ti 0.68Al 0.32B 1.35 thin films air-annealed at 700 °C for 1.0 h are shown in Fig. 11. After air-annealing, the total thickness of TiB 2.4 and Ti 0.68Al 0.32B 1.35 films increases by 14% and 7%, respectively. Fig. 11(a) and 11(b) show that TiB 2.4 has a significantly thicker oxide scale than Ti 0.68Al 0.32B 1.35. The oxidized layers formed on these films exhibit two different nanostructures. The oxide scale of TiB 2.4, inset in Fig. 11(a), has two regions; (i) an outer layer, ~90- nm thick, which is mostly composed of sub-micrometer equiaxed crystallites with an average size of ~100 nm, see supplementary Fig. S2(b), and (ii) an inner layer, ~420-nm thick, which exhibits a
Fig. 10. XPS elemental concentration depth profiles of (a) TiB 2.4 and (b) Ti 0.68 Al 0.32 B 1.35 thin films air-annealed at T a = 700 °C for t a = 1.0 h as a function of sputtering depth d . XPS values are normalized to the ToF-ERDA compositions.
columnar structure. This columnar layer consists of small columns formed near the oxide/film interface and wide columns elongated along the growth direction, in which the column width increases toward the top surface (V-shaped structure). Contrary to the ox- idized TiB 2.4, the oxide scale of the Ti 0.68Al 0.32B 1.35 films, ~300- nm thick, exhibits a featureless cross-sectional structure, inset in Fig. 11(b). This scale is composed of two sublayers in which the outer layer, ~130-nm thick, is not as compact as the ~170-nm inner layer. Both scales show insufficient adhesion to the unoxidized lay- ers due to the large thermal expansion
α
mismatch between the oxide scales and the unoxidized layers (α
rutile−TiO2 = 7.14 × 10 −6
K −1and
α
TiB2 = 7.6–8.6 × 10 −6K −1[109]). SAED patterns obtained from the oxide scales, insets in Fig.11(a) and 11(b), are composed of weak diffraction signals arising from the tetragonal rutile-TiO 2 phase. The SAED pattern of oxidized Ti 0.68Al 0.32B 1.35 also consists of a diffuse ring that indicates the presence of an additional oxide phase which is amorphous.The Z-contrast plan-view images of the TiB 2.4 and Ti 0.68Al 0.32B 1.35 thin films air-annealed at 700 °C for 1.0 h, acquired from areas A1 and A2 in Fig. 11(a) and 11(b), are ex- hibited in Fig. 11(c) and 11(d), respectively. The nanostructure of air-annealed TiB 2.4 shows ~14-nm-wide columns with porous boundaries in area A1, while the column width is significantly increased to ~37 nm and large gaps are formed between the columns in area A2, Fig. 11(c). Compared to TiB 2.4, the Z-contrast plan-view images of air-annealed Ti 0.68Al 0.32B 1.35 reveal nodular grains for both areas A1 and A2, Fig. 11(d). The nanostructure of the inner sublayer appears more compact than the outer one, which is porous.
Fig. 11. Typical XSTEM images, with inset SAED patterns, from (a) TiB 2.4 and (b) Ti 0.68 Al 0.32 B 1.35 thin films air-annealed at T a = 700 °C for t a = 1.0 h. The SAED patterns are acquired from the area indicated by circles in (a) and (b). Plan-view STEM images, with inset O EDX maps, from (c) TiB 2.4 and (d) Ti 0.68 Al 0.32 B 1.35 thin films air-annealed at T a = 700 °C for t a = 1.0 h. Areas indicated by rectangular boxes (A1 and A2) in (a) and (b) show the regions where plan-view STEM images are acquired from. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
Fig. 12. The oxide scale thickness d ox of TiB 2.4 and Ti 0.68 Al 0.32 B 1.35 thin films plotted as a function of air-annealing (a) temperature T a for t a = 0.5 h and (b) time t a at T a = 700 °C.
Oxygen EDX maps acquired from areas A1 and A2, insets in Fig.11(c), show that the dark regions in the Z-contrast images of air-annealed TiB 2.4are voids which behave like wide channels con- tinuously transferring oxygen to the unoxidized regions. The size of these voids changes along the oxide scale; wider in area A2 than area A1. The O EDX maps in the insets of Fig.11(d) exhibit that the outer layer of the Ti 0.68Al 0.32B 1.35oxide scale is porous (area A2); however, there is an almost uniform oxygen distribution in the in- ner one (area A1) indicating a more compact area which does not have large voids. To compare all elemental distributions in the re- gions shown in Fig.11(c) and 11(d), the corresponding EDX maps are exhibited in supplementary Fig. S4.
Fig. 12(a) compares the XSEM-obtained thicknesses of oxide scales d oxformed on the TiB 2.4 and Ti 0.68Al 0.32B 1.35thin films after air-annealing at different T afor t a = 0.5 h. Onset temperatures for the oxide-scale formation on TiB 2.4 and Ti 0.68Al 0.32B 1.35are 400 °C and 600 °C, respectively. The oxide growth rates of both films ex- ponentially increase at T a> 700 °C. The XSEM images of the films air-annealed at T a = 500 °C and 800 °C are compared in Fig.13. At T a = 500 °C, an ~290-nm oxide scale forms on the TiB 2.4 films, Fig. 13(a), while no detectable oxidation product can be found in the XSEM image of Ti 0.68Al 0.32B 1.35, Fig. 13(c). Fig. 13(b) exhibits that TiB 2.4 is completely oxidized after air-annealing at 800 °C (d ox = ~1900 nm), but only ~29% of Ti 0.68Al 0.32B 1.35is evolved into
an oxidized layer (d ox = ~470 nm) at this temperature, Fig.13(d). It indicates that alloying with Al significantly enhances the high- temperature oxidation resistance of TiB 2-rich Ti 1-xAl xB ythin films. The variations in d oxas a function of t afor T a = 700 °C, the tem- perature at which the oxidation product of TiB 2.4 contains ~0 at.% B, are plotted in Fig. 12(b). The TiB 2.4 thin films have thicker ox- ide scales than Ti 0.68Al 0.32B 1.35, with a thickness difference which becomes more pronounced for t a> 1.0 h.
4. Discussion
Adding Al to the sputter-deposited TiB y films via HiPIMS- assisted ion subplantation leads to changes in nanostructure. Both as-deposited TiB 2.4and Ti 0.68Al 0.32B 1.35thin films have a hexagonal columnar structure, while the crystallinity decreases significantly with a change in the crystal orientation from dominant 001 for TiB 2.4 to a mixed orientation of 001 and 101 for Ti 0.68Al 0.32B 1.35. Plan-view Z-contrast images, together with EDX maps and EELS spectra, show that the TiB 2.4 films are composed of nanocrystalline columns separated by a B-rich tissue phase, which is typical for sputter-deposited diboride films [94]. The Ti 0.68Al 0.32B 1.35 alloys, however, have Ti-rich columns surrounded by an Al-rich Ti 1-xAl xB y tissue phase which is B deficient.
A combination of XRD, XPS, SEM, STEM, SAED, and EDX anal- yses reveals that the sputter-deposited TiB 2.4 thin films have poor high-temperature oxidation resistance. For t a = 0.5 h, the thick- ness of the oxide scale formed on TiB 2.4 after air-annealing at T a = 500 °C is d ox = ~290 nm, while the films are completely oxidized at T a = 800 °C (d ox = ~1900 nm), Fig. 13(a) and 13(b). XPS and ToF-ERDA depth profiles in Figs.10(a) and supplementary S3(d), together with XRD results in Fig.7(a), show that the oxida- tion products of TiB 2.4 air-annealed at T a = 700 °C for t a = 1.0 h do not contain B and mostly consist of a tetragonal rutile-TiO 2 ( s) phase. Contrary to the bulk TiB 2synthesized by powder metallurgy processes [11, 52, 53], the as-deposited TiB y thin films grown by DCMS are typically overstoichiometric ( y > 2) and have colum- nar nanostructure in which the excess B segregates to the column boundaries and forms an amorphous B-rich tissue phase, shown in Figs. 3(c), 4(a), 4(b) and schematically depicted in Fig.14(a). At high temperatures in air, the B-rich column boundaries are highly prone to the formation of a B 2O 3( g) phase, with a vapor pressure that increases as a function of B concentration [56]. As a result of B 2O 3 ( g) evaporation, large gaps form between the TiO 2 columns, evident in Fig. 11(c), which act as wide channels for oxygen to readily access the unoxidized regions, causing a continuous vigor- ous oxidation. The oxidation-rate-limiting step is oxygen reaction at the oxide/film interface. The schematic of this process is illus- trated in Fig.14(b). This mechanism underlies the poor oxidation resistance of the sputter-deposited TiB 2.4 films with a B-rich net- work of column boundaries.
The oxide scale formed on TiB 2.4 at T a = 700 °C, schematically illustrated in Fig.14(c), is composed of two distinct TiO 2sublayers; the outer layer consisting of sub-micrometer equiaxed crystallites with an average size of ~100 nm, and the inner layer with a colum- nar structure. This columnar sublayer has small columns formed near the oxide/film interface and wide columns extended along the 001 direction, with open boundaries. The width of the TiO 2 columns increases toward the surface. The fine TiO 2 crystallites formed at the oxidation front undergo Ostwald ripening via the diffusive transfer of material, which depends on time and tempera- ture [110]. Thus, we observe that the TiO 2columns coarsen toward the scale surface creating an appearance of V-shaped columns as depicted in Fig.14(c). The coarsening of the TiO 2 columns is ac- companied by enlarging the porosities between the columns due to decreasing the surface to volume ratio of the columns. Therefore,
Fig. 13. Typical XSEM images of TiB 2.4 and Ti 0.68 Al 0.32 B 1.35 thin films air-annealed for t a = 0.5 h at T a = (a and c) 500 °C and (b and d) 800 °C.
Fig. 14. Schematic cross-sectional illustration of (a) as-deposited TiB 2.4 thin films, together with the nanostructure of TiB 2.4 thin films (b) during and (c) after oxidation experiment at 700 °C. The columns in (b) are shown at a higher magnification than (a).
the TiO 2 oxide scale does not provide a good oxidation protection of the underlying TiB 2.4 film.
Compared to TiB 2.4, the Ti 0.68Al 0.32B 1.35 alloy films exhibit sig- nificantly better high-temperature oxidation resistance. The XSEM images in Fig.13(c) and 13(d) indicate that there is no detectable oxide layer on Ti 0.68Al 0.32B 1.35 after air-annealing at T a = 500 °C for t a = 0.5 h, while the thickness of the oxide scale formed at T a = 800 °C is d ox = ~470 nm, which is considerably less than that of the TiB 2.4 scale (d ox = ~1900 nm) shown in Fig. 13(b). There can be different mechanisms responsible for this improvement. Al has a strong affinity to oxygen with an Al 2O 3 formation en- thalpy
H 25f ◦C (Al 2O 3) of −17.4 eV/atom [111], higher than that of Ti (
H 25f ◦C(rutile-TiO 2) =−9.8 eV/atom) [111], which results in the formation of an Al-containing oxide scale on the Ti 0.68Al 0.32B 1.35 alloys. This scale that suppresses bulk and surface diffusion can ef- fectively decrease the oxidation rate. The other effect can be at- tributed to the replacement of B by Al atoms in column bound- aries, which together with the formation of the Al-containing ox- ide scale, do not allow the TiO 2 columns to coarsen and form open column boundaries. The evidence for this scenario is the ab-
sence of X-ray reflections from the Ti 0.68Al 0.32B 1.35oxidation prod- ucts, Fig.7(b), which is due to the formation of amorphous phases. While the oxidation products of TiB 2.4 are highly porous and con- tain ~0 at.% B at T a = 700 °C, the air-annealed Ti 0.68Al 0.32B 1.35al- loys have much denser oxide scales with ~10 at.% B, Figs. 10 and 11.
Contrary to air-annealed Ti 1-xAl xN films with x≤ 0.64, in which the oxide scales consist of two sublayers – the outer Al-rich and the inner Ti-rich [107, 112], the XPS and ToF-ERDA depth pro- files acquired from the Ti 0.68Al 0.32B 1.35 alloy films air-annealed at T a = 700 °C for t a = 1.0 h do not show a significant change in the Al concentration profile as a function of depth, see Fig.10(b) and supplementary S3(h). The same trend was previously reported for Ti 1-xAl xN with x > 0.64 [106]. The oxide scale of Ti 0.68Al 0.32B 1.35 has a featureless cross-sectional structure composed of two sub- layers in which the outer layer (~130 nm) contains some porosities, while the inner layer (~170 nm) appears dense, inset in Fig.11(b). Similar to the air-annealed TiB 2.4 layers, the formation of these porosities is due to the coarsening of the oxide crystallites; how- ever, the presence of Al significantly suppresses the coarsening rate
and decreases the oxygen diffusion. Hence, Al addition together with the formation of a nanostructure with Al-rich column bound- aries result in a dense oxide scale which enhances the oxidation resistance.
5. Conclusions
We study the influence of Al addition on the oxidation proper- ties of sputter-deposited TiB 2-rich Ti 1-xAl xB y thin films. TiB 2.4 lay- ers are grown by DCMS from a TiB 2 target at a TiB 2-target power P TiB2 = 400 0 W and a negative dc substrate bias of 200 V, while
the Ti 1-xAl xB yalloy films are deposited by hybrid Al-HiPIMS/TiB 2- DCMS co-sputtering with a 200-V negative substrate bias synchro- nized to the Al-rich portion of each HiPIMS pulse. The compo- sition of as-deposited Ti 1-xAl xB y alloys are varied by increasing P TiB2, while all deposition parameters are maintained constant. The Al/(Ti +Al) ratio, x, decreases from x= 0.35 to 0.29 as P TiB2 is in- creased from 30 0 0 to 50 0 0 W, whereas the B/(Ti +Al) ratio, y, in- creases from 1.30 to 1.54. All as-deposited thin films show colum- nar structure. The column boundaries of TiB 2.4 films are B-rich, while the Ti 0.68Al 0.32B 1.35 alloys have Ti-rich columns surrounded by an Al-rich Ti 1-xAl xB y tissue phase which is B deficient. The sputter-deposited TiB 2.4 films exhibit rapid high-temperature ox- idation. The oxidation products of TiB 2.4 formed at temperatures T a > 500 °C do not contain B and mostly consist of a rutile-TiO 2 ( s) phase. The resulting oxide scales are highly porous due pri- marily to the evaporation of B 2O 3 ( g) phase as well as the coars- ening of TiO 2 crystallites. This poor oxidation is significantly im- proved by alloying with Al. While air-annealing at T a = 800 °C for t a = 0.5 h leads to the formation of an ~1900-nm oxide scale on TiB 2.4, the thickness of the Ti 0.68Al 0.32B 1.35 scale is ~470 nm. The enhanced oxidation resistance is mainly attributed to the for- mation of a dense, protective Al-containing oxide scale which decreases the oxygen diffusion rate by suppressing the oxide- crystallites coarsening. In addition to the improved oxidation re- sistance, the Ti 0.68Al 0.32B 1.35 alloy films with a nanostructure con- sisting of hard diboride-structure columns surrounded by Al-rich column boundaries exhibit low stresses, good indentation tough- ness, and maintain the high hardness of TiB 2.4 films.
Declaration of Competing Interest
The authors declare that they have no known competing finan- cial interests or personal relationships that could have appeared to influence the work reported in this paper.
Acknowledgments
We acknowledge support from the Knut and Alice Wallenberg (KAW) foundation for Project funding (KAW 2015.0043), a Fellow- ship/Scholar Grant, and support of the electron microscopy labo- ratory in Linköping. Financial support from the Swedish Research Council VR Grant 2014–5790, 2018–03957, and 642–2013–8020, the VINNOVA Grant 2018–04290, an ˚Aforsk foundation grant #16– 359, and Carl Tryggers Stiftelse contracts CTS 15:219, CTS 17:166, and CTS 14:431 are also gratefully acknowledged. Furthermore, the authors acknowledge financial support from the Swedish Govern- ment Strategic Research Area in Materials Science on Functional Materials at Linköping University (Faculty Grant SFO Mat LiU No. 20 09 0 0971). Supports from the Swedish research council VR-RFI (#2017–00646_9) for the Accelerator based ion-technological cen- ter and from the Swedish Foundation for Strategic Research (con- tract RIF14–0053; for the tandem accelerator laboratory in Uppsala University, and contract RIF14–0074; for the electron microscopy laboratory) are acknowledged.
Supplementary materials
Supplementary material associated with this article can be found, in the online version, at doi: 10.1016/j.actamat.2020.07.025. References
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