employing metal iodides and oxygen in ald and cvd of ...162869/...vi. growth of sno2 thin films by...

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Comprehensive Summaries of Uppsala Dissertations from the Faculty of Science and Technology 852 Employing Metal Iodides and Oxygen in ALD and CVD of Functional Metal Oxides BY JONAS SUNDQVIST ACTA UNIVERSITATIS UPSALIENSIS UPPSALA 2003

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  • Comprehensive Summaries of Uppsala Dissertationsfrom the Faculty of Science and Technology 852

    Employing Metal Iodides andOxygen in ALD and CVD of

    Functional Metal Oxides

    BY

    JONAS SUNDQVIST

    ACTA UNIVERSITATIS UPSALIENSISUPPSALA 2003

  • Till�Puman��från�Uman�

  • � � �

    List�of�papers��

    This�thesis�is�based�on�the�following�papers,�which�will�be�referred�to�in�the�text�by�their�roman�numerals.��

    I.� Atomic�layer�deposition�of�Ta2O5�by�using�the�TaI5�and�O2�precursor�combination�

    � J.�Sundqvist,�H.�Högberg,�A.�Hårsta�� Chem.�Vap.�Deposition�(Accepted)��

    II.� Atomic�layer�deposition�of�polycrystalline�HfO2�films�by�the�HfI4-O2�precursor�combination�

    � J.�Sundqvist,�A.�Hårsta,�J.�Aarik,�K.�Kukli,�A.�Aidla�� Thin�Solid�Films�427�(2003)�147��

    III.� Properties�of�hafnium�oxide�films�grown�by�atomic�layer�deposition�from�hafnium�tetraiodide�and�oxygen�

    � K.�Kukli,�M.�Ritala,�J.�Sundqvist,�J.�Aarik,�J.�Lu,�T.�Sajavaara,�M.�Leskelä,�A.�Hårsta�

    � J.�Appl.�Phys.�92�(2002)�5698��

    IV.� Atomic�layer�deposition�of�epitaxial�and�polycrystalline�SnO2�films�from�the�SnI4/O2�precursor�combination�

    � J.�Sundqvist,�A.�Tarre,�A.�Rosental,�A.�Hårsta�� Chem.�Vap.�Deposition�9�(2003)�21��

    V.� Chemical�vapour�deposition�of�epitaxial�SnO2�films�by�the�SnI4-O2�precursor�combination�

    � J.�Sundqvist,�M.�Ottosson,�A.�Hårsta�� Chem.�Vap.�Deposition�(submitted)��

    VI.� Growth�of�SnO2�thin�films�by�ALD�and�CVD:�a�comparative�study�� J.�Sundqvist,�A.�Hårsta�� Proceedings�of�the�sixteenth�Int.�CVD�Conf.,�Vol.�1�(2003)�511.��

    VII.� Microstructure�characterisation�of�ALD�grown�epitaxial�SnO2�thin�films�by�ALD�

    � J.�Lu,�J.�Sundqvist,�M.�Ottosson,�A.�Tarre,�A.�Rosental,�J.�Aarik,�A.�Hårsta�� (Manuscript)��

    VIII.� Nanoepitaxy�of�SnO2�on�αααα-Al2O3(0�1�2)�� A.�Tarre,�A.�Rosental,�J.�Sundqvist,�A.�Hårsta,�T.�Uustare,�V.�Sammelselg�� Surf.�Sci.�(In�Press)�

    iv�

  • � � �

    IX.� Gas�sensing�properties�of�epitaxial�SnO2�thin�films�prepared�by�atomic�layer�deposition�

    � A.�Rosental,�A.�Tarre,�A.�Gerst,�J.�Sundqvist,�A.�Hårsta,�A.Aidla,�J.�Aarik,�V.�Sammelselg,�T.�Uustare�

    � Sens.�Actuators�B�(In�press)��

    X.� ALD�of�some�metal�oxides�using�the�precursor�combination�iodide�and�oxygen�

    � J.�Sundqvist,�M.�Schuisky,�A.�Hårsta�� Proceedings�of�The�4th�International�Conference�on�Materials�for�

    Microelectronics�and�Nanoengineering�(2002)�65�

    Abbreviations��

    ALCVD� Atomic�Layer�Chemical�Vapor�Deposition�ALD� Atomic�Layer�Deposition�ALE� Atomic�Layer�Epitaxy�APB� Anti�Phase�Boundary�CVD� Chemical�Vapour�Deposition�DRAM� Dynamic�Random�Access�Memory�EDX� Energy�Dispersive�X-ray�spectroscopy�GI-XRD� Gracing�Incidence�X-Ray�Diffraction�HTEM� High�resolution�Transmission�Electron�Microscopy�MIM� Metal-Insulator-Metal�MOCVD� Metal�Organic�Chemical�Vapour�Deposition�MOSFET� Metal-Oxide-Semiconductor-Field-Effect-Transistor�PVD� Physical�Vapour�Deposition�QCM� Quartz�Crystal�Microbalance�QMS� Quadrupole�Mass�Spectrometer�RBS� Rutherford�Backscattering�Spectroscopy�RT� Room�Temperature�SAED� Selective�Area�Electron�Diffraction�SIMS� Secondary�Ion�Mass�Spectroscopy�TACVD� Thermally�Activated�Chemical�Vapour�Deposition�TEM� Transmission�Electron�Microscopy�TOF-ERDA� Time�Of�Flight�Elastic�Recoil�Detection�Analysis�XPS� X-ray�Photoelectron�Spectroscopy�XRD� X-Ray�Diffraction��XRFS� X-Ray�Fluorescense�Spectroscopy�XRR� X-Ray�Reflectivity�

    v�

  • � � �

    Contents�

    1� Introduction......................................................................................1�1.1� CVD..........................................................................................2�1.2� ALD..........................................................................................2�1.3� Precursor�selection.....................................................................5�

    1.3.1� The�metal�precursor...............................................................5�1.3.2� The�oxygen�precursor............................................................7�

    2� Experimental ..................................................................................11�2.1� Reactors ..................................................................................11�2.2� Characterisation�of�thin�films�-�analysis�techniques ..................13�

    3� ALD�of�high-k�oxides......................................................................17�3.1� Ta2O5 ......................................................................................18�

    3.1.1� Film�growth ........................................................................19�3.2� HfO2........................................................................................25�

    3.2.1� Film�growth ........................................................................25�3.2.2� Structure�and�composition ...................................................27�3.2.3� Electrical�characterisation....................................................30�3.2.4� Thermal�stability .................................................................34�

    4� ALD�and�CVD�of�SnO2...................................................................35�4.1� CVD�of�epitaxial�SnO2 ............................................................36�

    4.1.1� Growth�kinetics...................................................................36�4.1.2� Epitaxial�relationships .........................................................39�

    4.2� ALD�of�polycrystalline�and�epitaxial�SnO2...............................43�4.3� Comparison�with�between�the�SnI4�and�SnCl4�ALD-processes .49�

    4.3.1� Surface�morphology ............................................................50�4.3.2� Microstructure.....................................................................51�

    5� Concluding�remarks .......................................................................57�

    6� Acknowledgements .........................................................................59�

    7� References.......................................................................................61��

  • � � �

    1�

    1� Introduction�

    Thin�films�are�applied�on�surfaces�to�change�their�physical�and�chemical�properties.�They�are�an�integrated�part�of�almost�every�consumer�product�available�in�the�21:st�century� supermarket.� Thin� films�are� applied� to� make� a� surface� harder,� more� wear�resistant�and�more� lubricant� like� in�your�kitchen� frying�pan.�Thin� films�applied� to�circuit� boards� can� make� a� non-conductive� surface� conduct� electricity.� Corrosion�resistant�and�decorative� thin� film�coatings�are�applied�onto�outdoor�products.�Anti�reflective�coatings�are�applied�on�computer�and�TV�screens�and�eyeglasses�to�avoid�light�reflections.�Applied�to�parts�of�a�jet�engine,�thin�films�allow�the�airplane�to�use�less� fuel�and�by�using� thin� film�semiconductors,�smaller�electronic�devices� can�be�made.�These�are�just�some�of�the�examples�in�daily�life�where�thin�films�are�applied�to�improve�on�product�quality.�

    Metal� oxides� possess� a� variety� of� different� physical� and� chemical� properties.�They� can� be,� e.g.,� superconducting,� semiconducting,� ferroelectric,� antiferrielectric,�ferrielectric,� paraelectric,� pyroelectric,� piezoelectric,� ferromagnetic,� antiferri-magnetic,� paramagnetic,� ion� conductive,� dielectric,� gas� sensing,� catalytic,� or�chemically� inert.� Thin� films� with� any� of� these� properties� are� highly� interesting� to�apply�on�a�surface�or�in�combination�in�a�sandwich�structure�to�make�a�device�or�to�boost�the�performance�of�a�device.�

    There� are� various� ways� to� deposit� a� thin� film� of� a� material� on� an� arbitrary�substrate.�One�of�the�simplest�way�would�be�to�dissolve�the�desired�components�in�a�solvent�and�by�using�a�brush�apply�it�onto�the�substrate�and�then�let�it�dry�for�some�period�of�time,�something�close�to�a�Sol-gel�process.�Another�way�would�be�to�apply�brute� force� and� physically� deposit� the� material� onto� the� substrate,� like� in� Physical�Vapour�Deposition�(PVD).�A�more�elegant�way�is�to�let�surface�chemical�reactions�govern� the�growth�of� the�film� like� in�Chemical�Vapour�Deposition�(CVD)�or�even�more�elegantly�separate� these� reactions� in� time�and�digitally�control� the�growth�of�the�film�on�an�atomic� level,� like� in�Atomic�Layer�Deposition�(ALD).�All� thin�film�deposition� techniques� have� their� own� strong� points� and� disadvantages.� In� this�chapter� CVD�and� ALD� will� be� briefly� described� together�with�available� materials�from�which�the�films�can�be�deposited,�namely�the�precursors.�

  • � � �

    2�

    1.1� CVD�Excellent� textbooks�and�reviews�about� the� fundamentals�of�CVD�and�examples�of�various� processes� are� readily� available� today� [1-6].� The� general� concepts� of� CVD�are�here�briefly�described.��

    In�CVD�the�film�growth�is�governed�by�chemical�reactions�and�the�growth�rate�of�the� film� is�hence� controlled� by� the� rate� of� these� reactions.� There�are� a�number� of�ways�by�which�the�rate�can�be�controlled.�First�of�all�the�chemical�reaction�must�be�activated.�This�can�be�achieved�by� thermal�energy,�photons,�electrons�or� ions�or� in�combination.� The� results� in� this� thesis� are� based� on� thermally� activated� chemical�vapour� deposition� (TACVD)� and� the� thermal� energy� is� applied� from� the� furnaces�surrounding� the� reactor� tube� (Hot-wall� CVD).� When� the� temperature� is� increased,�atoms� and� molecules� will� move� faster� and� a� CVD-process� can� be� described� as�kinetically� controlled� if� the� growth� rate� of� the� film� exponentially� depends� on� the�temperature�or�more�precise,�if�a�straight�line�can�be�fitted�when�the�logarithm�of�the�growth�rate� is�plotted�as�a�function�of� the� inverse� temperature�–�an�Arrhenius�plot.�The�apparent�activation�energy�can�then�be�derived�from�the�slope�of�the�line.��

    If� the�growth�rate� increases�exponentially�with� temperature�there�is�a�possibility�that�at�some� temperature� the�growth� rate� is�so�high� that� the�supply�of� reactants� is�slower� than� the� growth� rate� of� the� film.� Then� the� process� is� in� mass� transport�control,� i.e.,� the� feed� rate� of� one� or� more� precursors� (reactants)� determines� the�growth�rate.�

    The� growth� rate� may� also� be� controlled� by� thermodynamics,� nucleation� or�homogenous� reactions.� There� are� various� parameters� that� can� be� optimised� to�achieve� different� types� of� controls� in� a� TACVD-process,� e.g.,� gas� flows,� type� of�flow�(turbulent�or�viscous),�the�feed�rate�of�the�precursors,�operating�pressure,�and�as�mentioned�the�temperature.�

    When� the� CVD-process� is� kinetically� controlled� a� film� with� homogenous�thickness� and� stoichiometry� can� be� deposited� on� complex� substrates� with� high�aspect� ratio� geometries� as� long�as� temperature� gradients� are� avoided.� Under� these�conditions� CVD� is� well� suitable� for� mass� production� of� thin� film� coatings� on� an�industrial�scale.��

    1.2� ALD�ALD,�also�denoted�ALE�or�ALCVD,�was� invented�by�T.�Suntola� in� the�1970’s�for�fabrication� of� polycrystalline� luminescent� ZnS:Mn� and� amorphous� Al2O3� insulator�films�for�electroluminescent�flat�panel�displays�[7].�

  • � � �

    3�

    Thin� film� growth� by� ALD� occurs� through� alternate� pulsing� of� the� precursor�fluxes�onto�the�substrate�surface�and�subsequently�the�precursor�species�chemisorbs�or�undergo�a�surface� reaction� [8].�The�reactor� is�purged�with�an� inert�gas�between�the�precursor�pulses�in�order�to�remove�the�by-products�and/or�remains�of�precursor�species.�By�careful�optimisation�of�the�experimental�conditions�the�process�proceeds�via� saturative� steps.� Under� these� conditions� the� growth� is� stable� and� the� thickness�increase� is�constant� in�each�deposition�cycle.�The�self-limiting�growth�mechanism�enables� growth� of� conformal� thin� films� with� accurate� thickness� control� on� large�areas�[9].�Recent�reviews�concerning�some�different�areas�of�ALD�are:�catalyst�[10],�nanotechnology� [11],� precursors� [12],� and� electronic� and� optoelectronic� materials�[13,�14].�

    Figure�1.1�The�growth�rate�as�a� function�of� temperature�in�a�general�ALD-process�[15].� This� schematic� figure� shows� some� of� the� different� processes� that� limit� the�temperature�span�of�the�ALD-window.��

    The�ideal�ALD-process�should�have�a�temperature�window�where�the�growth�rate�is� independent�of� the�deposition� temperature� -� the�ALD-window.�The� temperature�width�of� the�window�depends�on� four�different�surface�processes� [15].� Initially,�at�lower� temperatures� the�growth�rate� is�governed�by� either� thermal�activation�or�by�precursor� condensation� (or� multilayer� adsorption)� before� it� reaches� a� saturated�growth�rate�[16].�At�higher�temperatures� the�growth�rate�can�either� increase�due� to�precursor� decomposition� or� decrease� due� to� desorption� of� the� precursor.� Typical�upper� limiting�factors�are�desorption�of� the�metal�precursor�and/or�dehydroxylation�of� the� surface� with� increasing� temperature� in� H2O-based� ALD-processes� when�

    therm

    al�ac

    tivat

    ion

    condensationde

    com

    posit

    ion

    desorption

    ALD�window

    Gra

    owth

    �R

    ate

    Growth�Temperature

  • � � �

    4�

    growing� metal� oxides� [17].� The� dehydroxylation� takes� place� when� surface� –OH�groups�combine�and�liberate�from�the�surface�as�H2O(g):��

    2-OH���-O+�H2O(g)� � � (1.1)��

    Obviously,� additional� processes� may� control� the� growth� rate� even� though� the�process� is� performed� inside� the� ALD-processing� window� where� it� is� self-limiting.�For� instance� the� surface� morphology,� texture� and� phase� of� the� growing� film� may�affect�the�growth�rate�as�a�function�of�temperature,�especially�for�crystalline�growth.�Figure�1.2�shows� the�growth�rate�of� rutile�TiO2�and�SnO2� in�ALD� [18,�IV,�X].� In�both� cases� the� growth� rate� monotonically� increase� with� an� increase� in� deposition�temperature�in�the�case�of�polycrystalline�growth.�However,�when�epitaxial�growth�is� observed� the� growth� rate� saturates� and� only� shows� a� small� increase� as� the�temperature�is�elevated.��

    Figure� 1.2.� Growth� rate� of� SnO2� and� TiO2� films� grown� on� α-Al2O3(0�1�2)� and�Si(1�0�0)� substrates� as� a� function� of� deposition� temperature.� The� lines� are� to� be�considered�as�a�guide�for�the�eye.�

    Since� epitaxial� growth� is�not� realised� on� silicon� or� silica� substrates,� the�higher�growth� rate� observed� in� polycrystalline� growth� is� linked� to� a� rougher� surface�morphology�at�elevated�temperatures.�In�the�case�of�epitaxial�growth�there�is�reason�to�believe�that�the�number�of�available�adsorption�sites�is�better�defined,�whereas�on�roughening�surfaces� the�number�of�adsorption�sites� is�continuously� increasing.�The�small� increase� in� growth� rate� as� the� deposition� temperature� is� increased� could� be�connected� to�I2(g)�liberation�through�either�decomposition�at� the�surface�(1.2)�or�in�the�gas�phase�(1.3):�

    200 300 400 500 600 700 8000.00

    0.05

    0.10

    0.15

    0.20

    0.25

    0.30

    Epitaxial

    Poly�

    crysta

    lline

    Poly�

    crysta

    lline

    Epitaxial

    Gro

    wth

    �R

    ate�

    [nm

    /cyc

    le]

    Deposition�Temperature�[oC]

    �TiO2�on�Si(1�0�0)

    �TiO2�on�α-Al

    2O

    3(0�1�2)

    �SnO2�on�SiO

    2/Si(1�0�0)

    �SnO2�on�α-Al

    2O

    3(0�1�2)

  • � � �

    5�

    2MIx(ads)���2MIx-1(ads)�+�I2(g)� � (1.2)��

    MIx(g)���MIx-2(g)�+�I2(g)� � � (1.3)��

    In�both�cases� the�result� is� that� the�I2(g)� liberation� leaves�more�space�at� the�surface�and�additional�MIx(g)�may�adsorb�at�the�surface�when�the�deposition�temperature�is�increased.�It�is�worth�mentioning�that�process�(1.2)�was�observed�in�a�QCM�study�of�the�TiI4-O2�process� [19]�and�process� (1.3)� is�most�probably�present� in� the�SnI4-O2�process�according�to�thermodynamical�calculations�[V].�

    1.3� Precursor�selection�The� selection� of� precursors� is� as� important� in� CVD� as� it� is� in� ALD,� and� several�important� requirements� should� be� satisfied.� The� precursors� should� have� a� good�thermal� stability,� i.e.,� they� should� not� decompose� in� the� temperature� regime� used.�Furthermore,�they�should�easily�react�at�the�required�temperature,�if�possible�without�leaving�any�precursor�residues� in� the�film.�They�must�also�have�a�sufficiently�high�vapour�pressure�and�be�as�chemically�pure�as�possible.�However,�in�an�ALD-process�less� amount� of� precursor� is� required� to� saturate� the� surface� and� hence� precursors�with� lower�vapour�pressures�can�be�employed.�More�detailed�account�of�precursor�requirements�can�be�found�elsewhere�[3,�12].�

    1.3.1� The�metal�precursor�The�metal�precursors�used�are�generally�either�metalorganic1�or�halide�compounds.�Metalorganic�precursors�are�widely�used� in�MOCVD�due� to�several�reasons.�There�is� an�advantage� to� have� precursors� with� a�high� vapour� pressure� at� RT� in� a�CVD-process�due�to�easy�precursor�handling,�the�possibility�to�use�cold-wall�reactors�and�to�supply�a�high�feed�rate�of�the�precursor�in�fast�CVD-processes.�Furthermore,�there�is� a� need� for� multi-component� precursors� to� deposit� multi-component� oxides� like�Pb(Zr,Ti)O3,� (Ba,Sr)TiO3,� (Na,K)NbO3� and� YBaCuO� from� a� single� source.�Precursors�with�high�vapour�pressures�at�RT�and/or�multi�component�compounds�are�not�always�commercially�available�but�can�in�some�cases�be�synthesised.�One�of�the�most�widely�used�metal�precursors�in�ALD�is�TMA�(trimethyl�aluminium),�which�is�commonly�used� to�deposit�Al2O3� in� the�combination�with�H2O� [15,�20].�The�main�drawbacks� reported� from� metalorganic� based� CVD/ALD� are� connected� to� carbon�contamination� and� thermal� decomposition� of� the� organic� ligands� at� elevated�temperatures.� The� thermal� decomposition� (or� pyrolysis)� is� also� one� of� the�

    ������������������������������

    1�The�distinction�between�metalorganic�and�organometallic�is�here�disregarded.�

  • � � �

    6�

    advantages�in�employing�metalorganic�precursors�since�the�thermal�threshold�of�the�process�is�lowered�and�hence�deposition�is�feasible�at�lower�deposition�temperatures�at�least�compared�to�metal�chlorides.��

    Metal� halides,� especially� chlorides,� are� commonly� used� as� metal� precursors� in�both� CVD� [3]� and� ALD� [12,�15].� However,� employing� metal� iodides� in� CVD�[21-23]�and�ALD�[I-IV,�X,�22,�23]�has�reduced�the�halide�contamination�level�in�the�films�compared�to�metal�chloride�based�processes.�This�can�perhaps�be�explained�by�the� thermal�stability�of� the�precursor.� In�Table�1.1,� the�enthalpies�of� formation� for�some�iodides�and�chlorides�are�shown.�The�general�trend�is�that�metal�iodides�have�a�less�negative�enthalpy�of� formation� than� the�corresponding�metal�chloride,� i.e.,�the�iodides� are� less� stable.� The�higher� contamination� level� observed� in� chloride-based�processes�might�thus�be�connected� to� the�difficulty� in�removing� the�rather�strongly�bonded�chlorine�atoms.�In�contrast,�it�is�much�easier�to�remove�the�iodine�ligands�in�metal� iodides,� accordingly� resulting� in� purer� films.� It� can� be� mentioned� that� the�oxychlorides� have� a� more� negative� enthalpy� of� formation� than� the� corresponding�oxyiodides.� This� might� explain� why� it� is� more� difficult� to� achieve� contamination�free�films�using�chlorides.�The�lower�stability�of�the�metal�iodides�can�also�explain�why�crystalline�films�are�often�obtained�at� lower� temperatures.�The�gain�in�energy�going�from� the�metal�iodide� to� the�metal�oxide� is� larger�than�when�going�from� the�metal� chloride� to� the� metal� oxide,� thus� providing� the� necessary� energy� for�crystallisation.�[21]�

    Table�1.1.�Enthalpy�of� formation� for�some�metal�chlorides�and� iodides�relevant� to�this�work.�

    � � � �

    MClx� ∆H0f,298/kJ.mol-1� MIx� ∆H0f,298/kJ.mol-1�TiCl4� -763�[24]� TiI4� -277�[24]�ZrCl4� -870�[24]� ZrI4� -362�[24]�HfCl4� -891�[25]� HfI4� -371�[26]�TaCl5� -765�[24]� TaI5� -266�[27]�SnCl4� -472�[28]� SnI4� -89�[29]�

    A� disadvantage� in� chloride-based� processes� is� the� etching� of� the� metal� oxide�deposit�by�the�metal�chloride�itself.�It�has,�e.g.,�been�shown�that�there�is�an�etching�process�present�in�the�TaCl5�based�processes�that�produces�volatile�oxychlorides�by�the�equilibrium�(1.4)�[30]:���

    Ta2O5(s)�+�3TaCl5(g)���5TaOCl3(g),�∆H0f,298�=�-892�kJ.mol-1� (1.4)��

  • � � �

    7�

    The� etching� process� has� been� described� in� detail� in� the� case� of� ALD� of� Ta2O5�from�TaCl5�and�H2O�by�in-situ�QCM�and�film�growth�[31,�32].�It�was�concluded�that�due�to�etching�of�Ta2O5�by�TaCl5�via�volatile�oxychlorides�the�deposition�parameters�had� to� be� carefully� tuned� with� respect� to� the� TaCl5� pulse� length� and� deposition�temperature� to� obtain� uniform� films� with� an�acceptable� growth� rate.� Furthermore,�since�H2O�(or�H2O2)�has�to�be�used�when�employing�metal�chlorides�to�remove�the�chlorine�as�HCl�(see�next�section)�further�etching�of� the�metal�oxide�film�might�be�present.�

    1.3.2� The�oxygen�precursor�The� oxygen� precursor� is� in� most� cases� H2O� but� H2O2,� N2O,� O3� or� O2� can�also� be�employed� [12].� The� oxygen� can� also� be� incorporated� in� the� metal� precursor�as� in�metal�alkoxides�[33].�In�ALD-processes�to�deposit�rare�earth�oxides�O3�is�needed�to�activate�the�process�[34,�35].�

    The�classical�ALD-process� to�grow�metal�oxides� is�usually�based�on�employing�H2O� as� oxygen� precursor.� It� has� been� shown� that� H2O� is� excellent� for� growing�amorphous�films.�The�growth�mechanism� is� in� this�case�based�on�surface�hydroxyl�groups,�and�can�be�written�as�(exemplified�by�the�TiCl4-H2O�process�[36,�37]):��

    2�-OH(ad)�+�TiCl4(g)���(-O-)2TiCl2(ad)�+�2HCl(g)� (1.5)��

    (-O-)2TiCl2(ad)�+�2H2O(g)���(-O-)2Ti(OH)2(ad)�+�2HCl(g)� (1.6)��

    In� the� growth�mechanism�above,� studied� in-situ�at� 150�°C� by� QCM� and� QMS�[37],� TiCl4� reacts� with� the� surface� -OH� groups� and� HCl� is� released.� During� the�subsequent�H2O�pulse�the�two�remaining�surface�-Cl�groups�react�with�H2O�forming�volatile� HCl� and� surface� -OH� groups.� At� higher� temperatures� (above� 350�°C)� the�reaction�mechanism�is�changing�to�a�path�where�the�incoming�metal�precursor�reacts�with� only� one� surface� -OH� group.� It� has� also� been� shown� that� the� partial� H2O�pressure�can�control�the�film�structure�in�the�TiCl4-H2O�process�at�400�°C�[38].�

    If� as-deposited� crystalline� films� are� required,� either� polycrystalline,� textured� or�epitaxial,�a�higher�deposition� temperature� is�normally� required.�In� this�connection,�the� growth� rate� decreases� in� water� and� metal� chloride� based� ALD-processes� at�higher�temperatures�due�to�decreased�surface�density�of�active�adsorption�sites,�i.e.,�–OH�groups.�In�contrast,�if�using�metal�precursors�that�are�reactive�with�oxygen,�the�process� would�not�be� limited�by�surface�dehydroxylation�at�elevated� temperatures.�The�TiI4-O2�ALD-process�has�been�studied�by�QCM�[19]�and�the�following�growth�mechanism�was�suggested:��

    TiI4(g)���TiIx(ad)�+�(4-x)/2I2(g)� � (1.7)�

  • � � �

    8�

    TiIx(ad)�+�O2(g)���Ti(-O-)2(ad)�+�x/2I2(g)� � (1.8)��

    In�general�it�can�be�noted�that�it�is�important�in�any�metal�oxide�ALD�process�to�optimise�the�process�not�only�with�respect�to�the�oxygen�precursor�pulse�length�but�also� to� the� partial� pressure� of� the� oxygen� precursor.� If� the� ligands� bonded� to� the�metal�atoms�are�not� fully� removed�after� the�oxygen�precursor�pulse� they�might�be�incorporated�into�the�film�and�stay�as�a�contamination�that�most�probably�will�affect�the�film�quality� in�a�negative�way.�The� importance�of� the�oxygen�precursor�partial�pressure� has� been� surveyed� for� a� number� of� metal� oxide� ALD-processes.� As�mentioned�earlier,�the�partial�H2O�pressure�can�affect�the�structure�of�the�film�[38].�In� the�case�of�a�H2O�based�process� it�was�found�that�by� fully�optimising�the�ALD-process,�with�respect� to� the�partial�H2O�pressure�during� the�H2O�pulse,� the�growth�rate�for�a�series�of�metal�oxides�could�be�increased�[39].��

    Table�1.2.�Growth�rates�for�some�selected�metal�oxide�ALD-processes.��Metal�precursor�

    Oxygen�precursor�

    Tgrowth�[°C]�

    Growth�rate�[Å/cycle]�

    Crystallinity*� Ref.�

    Ta(OEt)5� H2O� 300� 0.49� am� [39,�40]�TaCl5� H2O� 300� 0.66� am� [39,�41]�TaI5� H2O-H2O2� 300� 0.75� am� [42]�TaI5� O2� 600� 1.7� poly� [I]�HfCl4� H2O� 300� 0.5� poly� [43]�HfI4� H2O� 300� 0.6� poly� [43]�HfI4� H2O-H2O2� 300� 0.75� poly� [44]�HfI4� O2� 500� 1.1� poly� [II]�Ti(OEt)4� H2O� 300� 0.6� poly� [39,�45]�Ti(OiPr)4� H2O� 300� 0.6� poly� [39,�46]�TiCl4� H2O� 300� 0.5� poly� [39,�47]�TiI4� H2O-H2O2� 300� 0.5� epi� [48]�TiI4� O2� 300� 0.9� epi� [18]�SnCl4� H2O� 300/600� 0.3/0.25� tex/epi� [50,�VII]�SnCl4� H2O-H2O2� 300/600� 0.3/0.25� tex/epi� [49,�50,�VII]�SnI4� H2O-H2O2� 300/600� 0.25/0.4� epi/epi� [50,�VII]�SnI4� O2� 600� 1.1� epi� [IV,�VII]�

    *�am�=�amorphous,�poly�=polycrystalline,�tex�=�textured�and�epi�=�epitaxial�

    In�Table�1.2�the�growth�rates�are�compared�between�different�metal�oxide�ALD-processes� relevant� to� this�work.�A�striking�difference� in� this�and� related�work�has�been� that� many� of� the� metal� iodide� based� ALD-processes� have� reached� a� higher�growth� rate� compared� to� the�metal� chloride� based� ALD-processes.� The� reason� for�

  • � � �

    9�

    this�observation�has�been�discussed�earlier�and�is�almost�certainly�due�to�etching�by�HCl� that� is� released� during� growth� (reaction� 1.5� and� 1.6)� and/or� by� the� metal�chloride�itself�as�in,�e.g.,�the�TaCl5-H2O�process�(1.4).�

    When�comparing�the�O2�and�–OH�based�metal�oxide�processes�it�is�clear�that�the�growth�rates�are�higher�for�the�former.�The�main�reason�for�this�observation�is�most�probably� that� the� film� depositions� were� performed� at� higher� temperatures.� If� the�same� deposition� temperature� where� to� be� used� in�a� –OH� based� process,� increased�surface�dehydroxylation�would�explain�the�lower�growth�rate�observed.�Although�O2�based� ALD-processes� are� most� commonly� realised� with� metal� iodides,� there� is� a�recent� report� employing� the� ZrCl4-O2� precursor� combination� [51].� If� comparing�metal�chloride�and�metal� iodide�based�process� in�general�at�elevated� temperatures,�the�MIx�species�adsorbed�at�the�surface�may�liberate�additional�iodine�through�I2(g)�release�due�to�the�relatively�lower�stability�of�metal�iodides�and�therefore�leave�more�space� for�additional�adsorption�of�MIx�species� resulting� in�a�higher�growth�rate.�A�further�increase�in�growth�rate�at�elevated�temperatures�might�also�be�due�to�surface�roughening,� which� was� seen� in� the� SnI4-O2� process� [IV]� when� comparing�polycrystalline�and�epitaxial�growth,�i.e.,�3D�compared�to�2D�growth.�

  • � � �

    10�

  • � � �

    11�

    2� Experimental�

    All� thin� film� depositions� reported� in� this� thesis� have� been� made� in� two� different�ALD� reactors.� The� reactors� are� homebuilt� research� reactors� and� have� their� own�drawbacks�and�restrictions.�However,�working�with�these�kinds�of�reactors�also�have�some� advantages� since� they� can� be� easily� modified� to� meet� the� demands� of� new�ideas.�

    2.1� Reactors�Both�of�the�reactors�are�flow�type�hot�wall�reactors�with�quartz�reactor�tubes,�which�allow�depositions�at�temperatures�higher�than�generally�employed�in�ALD.���

    Figure� 2.1.�A� schematic�drawing� of� ALD� reactor� 1.� The� numbers� indicated� in� the�figure�relates�to�the�gas�pulsing�sequence�that�is�described�in�table�2.1.�

    Reactor� 1� (Figure� 2.1)�has� four� furnaces� separately� monitored� by� temperature-control� units.� The� gas� pulsing� sequence� is� realised� by� computer� controlled�pneumatic�valves�by�closing�or�opening�a�pressured�air�line.�The�N2�and�O2�flows�are�controlled�by�mass�flow�controllers.�An�ALD�cycle�is�obtained�by�switching�on�and�off� the�gas�flows�as�summarised� in�Table�2.1.�The�solid�metal� iodide� is�evaporated�from� a� quartz�ampoule� with�a� 5� x� 10�mm� opening� that� is� located� in� tube��.�The�amount� of� metal� iodide� was� carefully� weighed� before� and� after� each� deposition�

    Heater�coils

    ����

    �� Exhaust�

    Substrate�holder�

    Substrates�

    ��

    ��

    ��

  • � � �

    12�

    experiment.�As�a�standard�procedure�no�more�than�50�w%�was�allowed�to�evaporate�during�an�experiment.�By�carefully�monitoring�the�mass�loss,�any�changes�compared�to�other�similar�experiments�are�a�strong�indication�of�leaks�or�possible�ageing�of�the�precursor.� In� the� CVD� experiments� [V,� VI]� both� precursor� fluxes� were� supplied�simultaneously.��

    Table�2.1.�The�ALD-pulsing�sequence�employed�in�reactor�1.�The�numbers�indicated�in�the�third�column�relates�to�the�numbers�in�Figure�2.1.��

    � Pulse� Tube�Flows�t1� MIx�Pulse� ��N2�carrier�gas�and�5%�back�flow�

    ��Closed���N2�bulk�flow�

    t2� Purge�1� ��5%�back�flow���N2�protective�flow���N2�bulk�flow�

    t3� O2�Pulse� ��5%�back�flow���N2�protective�flow���O2�flow�

    t4� Purge�2� ��5%�back�flow���N2�protective�flow���N2�bulk�flow�

    Reactor�2�(Figure�2.2)�has�a�similar�construction�and�is�described�in�more�detail�elsewhere� [52].�The�major�difference� is� that� it�has�a�more�compact�design�and� the�reactants�are�all�introduced�in�front�of�the�substrate�holder�through�a�central�orifice.�The�more�compact�design�allows�shorter�pulse�times�to�be�used.�The�reactor�is�also�equipped� with� a� mass� sensor,� a� Quartz� Crystal� Microbalance� (QCM),� that� can� be�used� for� monitoring� the� growth� of� an� ALD-process.� However,� the� QCM� was� not�used� in� the�papers�summarised� in� this� thesis,�since�all� investigated�processes� were�possible�only�above�the�temperature�limit�of�this�crystal.�

  • � � �

    13�

    Figure�2.2.�Schematic�drawing�of�ALD�reactor�2.��

    The� operating� pressures,� flow� rates,� pulse� times� (t1-t2-t3-t4),� deposition�temperatures� and� evaporation� temperature� of� the� metal� iodides� are� reported�separately�in�each�process�paper�[I,�II,�IV�and�V].�

    2.2� Characterisation�of�thin�films�-�analysis�techniques�The� deposited� thin� films� have� been� characterised� with� respect� to� their�microstructure,� stoichiometry,� impurities,� surface� morphologies,� deposition� rates�(thickness),� dielectric� properties,� and� gas� sensing� properties� with� various� analysis�techniques.�Table�2.2.�summarises�the�analysis�techniques�employed�in�the�different�papers�included�in�this�thesis�and�the�main�information�obtained.�

    XRD�[53]�was�used�as�the�main�analysis�tool�for�all�the�deposited�films.�Standard�θ-2θ� scans� were� mainly� used� for� textured� films� and� GI-XRD� was� employed�whenever� the� films� were� predominately� non-textured.� The� main� advantage� of� GI-XRD� is� the� larger� sample� volume� analysed,� i.e.,� thinner� films� compared� to� θ-2θ�scans�can�be�analysed.�By�changing�the�angle�of�the�incident�beam�inhomogeneities�in� phase� composition� with� respect� to� depth� can� also� be� probed.� For� textured� and�epitaxial�films,�rocking�curves�of�peaks�from� the�main�growth�direction�of� the�film�were�examined�with�respect�to�the�full�width�at�half�maximum�(FWHM)�values�as�a�measure�of�how�well�the�film�was�aligned�parallel�to�the�substrate.�To�determine�the�

    N2�+�Reactants

    N2

    N2

    Exhaust

    Mass�sensorSubstrate

    ThermocoupleSubstrate�holder

    HeaterReactor�tubes

  • � � �

    14�

    in-planar�orientational�relationship�between�the�film�and�the�substrate,�ϕ-scans�from�a�film�and�a�substrate�reflection�that�is�not�from�a�plane�parallel�to�the�contact�plane�were� performed.� To� further� analyse� textured/epitaxial� films� pole� figures� were�recorded,�which�give�a�more�extensive�survey� of� the� film�by�detecting�minor� film�orientations�not�seen�in�the�ϕ-scans.�

    Table�3.2.�Analysis�techniques�used�for�characterisation.�

    Techique� Information� Paper�XRD�techniques� � �

    θ-2θ� Phase�content� IV-VI,�VIII,�IX�GI-XRD� Phase�content,�phase�

    inhomogeneity�I-V�

    Rocking�curve� Crystallite�alignment� IV-VI,�VIII,�IX�ϕ�scans� Epitaxial�relationships� IV-IX�Pole�figures� Film�orientations� V,�VII,�VIII�XRR� Thickness,�density,�

    roughness,�interface�I,�II,�IV,�V�

    AFM� Surface�morphology� IV,�VIII,�IX�TEM� Microstructure,�defects,�

    interface�III,�VII,�X�

    XPS� Impurities,�bonding�state� I,�V�RBS� Thickness,�interface,�

    stoichiometry�I�

    TOF-ERDA� Impurities� II,�III�XRFS� Relative�thickness,�

    impurities�I-VI,�VIII,�IX�

    XRR�was�used�mainly�to�determine�the�film�thickness�by�fitting�the�XRR-curves�with�help�of�the�software�WinGixa.�The�fitting�also�yields�other�useful�information�such� as� the� density,� surface� roughness� and� interface� roughness.� However,� these�additional�data�has�been�treated�with�caution.�

    XRFS�was�used�as�a�standard�analysis�tool�when�determining�the�growth�rate�in�a�deposition� process.� The� integrated� peak� area� of� the� metal� peak� can� be� linearly�related� to� the� film� thickness� if� assuming� no� deviation� in� density� of� the� deposited�films.� If� standards� with� known� thickness� are� available� this� is� a� very� fast� way� of�getting� the� film� thickness� and� the� growth� rate.�The� film� thickness� standards� were�made�by�depositing� the� films�by�ALD�and� then�calculate� the� thickness� from�XRR�data.� XRFS� can� also� give� an� estimate� of� any� impurities� present� in� the� film� for�

  • � � �

    15�

    elements�with�an�atomic�number�of�11�(Na)�and�above.�Since�the�detection�limit�of�iodine�is�below�1�at%�(according�to�the�manufacturer)�even�small�amounts�of�iodine�are� easily� detected.� Impurity� levels� (mainly� iodine),� and� in� some� cases� film�elemental� stoichiometry� (oxygen/metal� ratio),� have� also� been� analysed� by� TOF-ERDA�and�XPS.�In�contrast�to�XRFS,�these�two�techniques�also�give�the�elemental�composition�as�a�function�of�film�depth.�

    TEM� studies� were� used� to� investigate� the� microstructure� of� the� epitaxial� SnO2�films�(section�4.3.2.).�The�advantage�of�TEM�is�that�localised�information�from�the�film,� e.g.,� the� interface� between� the� film� and� substrate,� twin� and� antiphase�boundaries�can�be�obtained.�By�means�of�high�resolution�TEM,�dislocations�at� the�interface�and�at�twin�boundaries�can�be�observed.�The�epitaxial�relationship�between�the�film�and�substrate�can�be�investigated�by�SAED.�Additionally,�different�type�of�twins�can�be�distinguished�by�SAED.�These�results�may�also�act�as�a�support�for�the�results�from�XRD�investigations.��

    The�surface�morphology�of�some�deposited�films�has�also�been�studied.�In�these�studies� AFM� and� SEM� were� employed.� AFM,� as� compared� to� SEM,� can� give�quantification� in� the� lateral�dimension,�e.g.,� in� the� form�of� rms� (root-mean-square)�values�over�a�certain�scan�area,�whereas�SEM�only�gives�a�qualitative�picture.�

  • � � �

    16�

  • � � �

    17�

    3� ALD�of�high-k�oxides�

    Though� the� miniaturization� of� MOSFET� devices� has� offered� outstanding�improvement� in� performance,� it� seems� to� have� reached� a� fundamental� limit� with�SiO2� (εr� =� 3.9)� as� the� gate� dielectric.� For� SiO2� layers� about� 10�Å� or� thinner,� the�leakage� current� between� the� gate� and� the� substrate� poses� a� major� concern,�particularly�in�relation�to�the�standby�power.�For�example,�if�the�maximum�tolerable�gate� leakage�current� is�1�A/cm2,� the�oxide� thickness�may�not�be� thinner� than�14�Å�[54].� In� an� attempt� to� improve� on� these� limitations,� there� has� been� an� extensive�search� for� high-k� (εr� =� 9-25)� gate� dielectrics� to� replace� SiO2.� In� order� to� meet�performance� and� leakage� current� requirements� for� low� power� applications,�high-k�gate�dielectrics�are�required�by�the�year�2005�[55].�Among�the�high-k�dielectrics�that�are�widely�considered�as�promising�candidates�for� this�purpose�are�ZrO2�and�HfO2�[56,�57].� Mixed� oxides� and� silicates�are� also� gaining� interest� as�alternative� high-k�gate�dielectrics�[58,�59].�Ta2O5�was�previously�considered�as�a�candidate,�but�since�it�is�not�chemically�stable� on�silicon�and� its�band�gap� (4.4�eV)� is�not�symmetrically�aligned�with�silicon� leaving�only�a�0.3�eV� barrier�against�electron� injection� to� the�conduction�band�[60]�it�is�no�longer�considered�as�a�proper�replacement�for�SiO2�in�the�case�of�silicon-based�MOSFET�structure.��

    One�of� the�main�advantages� in�silicon�technologies�has�been� the�easy�formation�of�SiO2�on�the�silicon�surface,�which�up�to�now�has�worked�perfectly�as�a�dielectric�layer.�This�advantage�now�seems� to�be�a�disadvantage�due�to� the�fact�that� it�seems�impossible�to�grow�any�high-k�oxide�material�on�top�of�silicon�without�the�formation�of� an� interfacial� SiO2� layer� due� to� the� strong� affinity� of� silicon� towards� oxygen.�Either�the�SiO2� layer�forms� initially�during�the�growth�process�or� later�because� the�high-k�oxide�is�not�stable�on�silicon.�[33]�

    In� memory� applications� such� as� the� DRAM� the� constraints� are� somewhat�different� since� the� structure� can� be� based� on� a� metal� –� insulator� –� metal� stack�(MIM).� In� this� case� a� metal� that� has� a� work� function� that� is� higher� than� that� of�silicon,�such�as�ruthenium�and�platinum,�may�be�employed�and�hence� the�problem�with�electron�injection�into�the�conduction�band�can�be�solved.�Furthermore,�oxides�that�are�not�stable�on�silicon,�such�as�Ta2O5,�may�also�be�considered.�

    The�main�route�to�minimise�the�leakage�current�is�to�reduce�the�amount�of�grain�boundaries�in�the�high-k�oxide,�thus�amorphous�or�single�crystalline�epitaxial�films�are�preferred.� In� the�case�of� the�ALD-processes�studied� in� this� thesis,�TaI5-O2�and�

  • � � �

    18�

    HfI4-O2,�both�resulted� in�polycrystalline�film�growth.�However,�both�metal� iodides�can� be� used� to� grow� amorphous� films� in� the� as-deposited� state� by� applying� a�different�oxygen�precursor,�H2O�or�H2O2�[44,�43].�When�targeting�amorphous�films,�the�metal�iodide�processes�could�be�advantageous,�since�ultrathin�films�grown�from�metal� iodides� have� shown� a� larger� tendency� to� form� amorphous� films� than� films�grown�by�metal�chloride�processes�[43].�The�reason�why�polycrystalline�growth�was�observed� in� the� case� of� the� two� O2-based� ALD-processes� is� due� to� the� higher�temperatures� needed� for� thermal� activation� compared� to� the� corresponding� -OH�group�based�ALD-processes.�

    Even� if� a� perfectly� amorphous� film� could� be� deposited� without� any� interfacial�SiO2�layer�there�is�still�one�problem,�namely�that�it�has�to�withstand�a�last�annealing�step�(900-1200�°C)�for�dopant�activation�without�crystallisation,�which�has�not�yet�been�achieved.�A�recent�review�of�high-k�gate�dielectrics�can�be�found�in�ref.�[61].�

    As� mentioned� above,� polycrystalline� films� are� not� desirable� in� the� case� of�MOSFET�and�DRAM�applications.�It�is�therefore�important�to�study�the�structure�of�the� grown� films� and� to� evaluate� the� formation� conditions� and� dependence� on�processing� parameters� to� gain� understanding� of� the� general� and� fundamental�processes�governing� thin�film�growth�of� these�metal�oxides.�Furthermore,� there�are�potential�applications�outside�the�field�of�microelectronics�for�polycrystalline�films.�

    3.1� Ta2O5�Tantalum� oxide,� Ta2O5,� has� large� technological� interest,� especially� because� of� its�optical� and� electrical� material� properties.� Thin� films� of� tantalum� oxide� can� be�applied�as�optical�coatings�[62],�ion-sensitive�membranes� in�solid-state� ion�sensors�[63],�gate�dielectric� in�metal-oxide-semiconductor� transistors� [64],�barrier�material�in�magnetic�tunnel�junctions�[65],�corrosion�resistant�coatings�[66],�and�as�a�high-k�dielectric�material�in�dynamic�memory�capacitor�structures�[67].�Recently�the�TaI5-O2�ALD-process�was�integrated�to�deposit�a�MFIS�diode�consisting�of�the�following�stack�Au/Na0.5K0.5NbO3/Ta2O5/Si�[68].�In�earlier�studies�SiO2�was�used�as�insulating�layer�but� it� is�generally�known� that�alkali� ions�have�a�high�diffusion�rate� into�SiO2�and�will�hence�reduce� the�lifetime�of� the�device.�Therefore�Ta2O5�was�deposited� to�act�as�both� insulator�and�diffusion�barrier�and� the�performance�of� the�MFIS-diode�was�compared�to�that�of�the�stack�Au/Na0.5K0.5NbO3/SiO2/Si.�The�use�of�Ta2O5�was�found� to� reduce� the� flat-band� voltage� (VFB)� deviation� that� has� been� attributed� to�intermixing�of�Na�and�K� ions� into� the�SiO2� layer.�The�novel�MFIS-diode�was�also�found� to� exhibit� a� wide� memory� window� (4.5� V),� low� leakage� current� and� long�retention�times�at�zero�bias.�

  • � � �

    19�

    0 1 2 3 4 5 6

    0.00

    0.05

    0.10

    0.15

    0.20

    0.25

    Gro

    wth

    �R

    ate�

    [nm

    �/�c

    ycle

    ]

    t�[s]

    �TaI5�Pulse�Length,�t

    1-4-4-4�s

    �O2�Pulse�Length,�4-4-t

    3-4�s

    3.1.1� Film�growth�In� the�TaI5-O2�process� the�pulse� times� for� the�metal�and� the�oxygen�precursors� (t1�and� t3)�were�changed� individually� to�achieve�self-limited�growth� (Figure�3.1).�The�N2�purging�times,�t2�and�t4�were�held�constant�at�4�s�to�assure�complete�separation�of�the�precursor�pulses.�The�optimisation�of�the�ALD-process�was�performed�using�HF-stripped� Si(1�0�0)� substrates� held� at� a� temperature� of� 600�°C,� by� applying� 500�deposition�cycles�and� then�measuring� the� film� thickness� to�obtain� the�growth�rate.�As� can� be� seen� in� Figure� 3.1� the� TaI5� pulse� (t1)� and� the� O2� pulse� (t3)� reaches�saturation�at�about�the�same�pulse�times.�Since�no�etching�of�the�growing�Ta2O5�film�could�be�observed�when�applying�longer�TaI5�pulses,�the�pulse�times�were�set�to�4-4-4-4�s�(t1-t2-t3-t4)�in�all�subsequent�deposition�experiments.�

    Figure�3.1.�Growth�rate�of�Ta2O5�as�a�function�of�the�duration�of�the�TaI5-pulse,�t1,�and� O2-pulse,� t3,� performed� on� Si(1�0�0)� substrates� at� a� deposition� temperature� of�600�°C.���

    After�these�initial�studies�a�temperature�series�from�400�to�750�°C�was�deposited�on� Si(1�0�0)� substrates� to� study� the� growth� rate� dependence� on� deposition�temperature� (Figure�3.2).�At�400�and�450�°C�continuous�Ta2O5� films�could�not�be�produced� by� applying� 500� deposition� cycles.� However,� continuous� films� were�achieved�at�500�°C�and�up� to�700�°C.�The�growth�rate�was�found� to� increase�from�0.08�nm/cycle�at�500�°C� to�0.17�nm/cycle�at�600�°C.�An� increase�of� the�deposition�temperature� above� 625�°C� resulted� in� a� dramatically� decreased� growth� rate,�0.03�nm/cycle�at�700�°C.�According� to�a� thickness�series�grown�at�600�°C� the�film�thickness�was�proportional�to�the�number�of�deposition�cycles�applied.�

  • � � �

    20�

    400 450 500 550 600 650 700

    0.00

    0.05

    0.10

    0.15

    0.20

    0.25

    Continuous�films

    Gro

    wth

    �R

    ate�

    [nm

    �/�c

    ycle

    ]

    Deposition�Temperature�[oC]

    Figure� 3.2.� Growth� rate� of� Ta2O5� as� a� function� of� deposition� temperature� on�Si(1�0�0)�substrates�by�applying�500�deposition�cycles.��

    The� phase� analysis� by� GI-XRD� revealed� that� the� Ta2O5� films� were� fully�crystallised� starting� from� 500� °C� [I].� It� was� clear� from� the� XRD� patterns� that� the�films� consisted� of� the� orthorhombic� β-Ta2O5� phase,� which� agrees� well� with� other�studies,�by�both�CVD�[69]�and�ALD�[42,�70],�of�Ta2O5�thin�films�grown�on�the�same�kind�of�substrates.�

    The� evolution� of� the� film� growth� was� studied� at� 475� °C� by� applying� different�number�of�deposition�cycles.�This�temperature�lies�between�the�temperatures�where�non�continuous�and�continuous�films�may�be�deposited�by�applying�500�deposition�cycles�and�is�hence�interesting�for�this�type�of�study.�The�SEM�pictures�(Figure�3.3)�can�be�regarded�as�snap-shots�of�the�film�growth�up�to�1000�deposition�cycles.�It�can�be� seen� that� grains� of� Ta2O5� nucleates� randomly� but� evenly� on� the� surface.� The�grains� then�continue� to�grow�and�new�grains�continuously�nucleate� in�between� the�primary�grains.�After�1000�deposition�cycles�something�close�to�a�continuous�film�is�finally�achieved�but�voids�are�still�present�between�the�grains.�

  • � � �

    21�

    a)� b)�

    d)�c)�

    1�µµµµm�

    1�µµµµm�

    a)� b)� c)�

    Figure�3.3.�SEM�pictures�of�depositions�performed�at�475�°C�by�applying�different�number�of�deposition�cycles:�(a)�200,�(b)�400,�(c)�500�and�(d)�1000�cycles.��

    The�earlier�mentioned� thickness�series�deposited�at�600�°C,�was�also�studied�by�SEM� (Figure�3.4).�At� this�deposition� temperature� it�can�be�seen� that� the� films�are�continuous�all�the�way�from�300�deposition�cycles.��

    Figure�3.4�SEM�pictures�of�depositions�performed�at�600�°C�by�applying�different�number�of�deposition�cycles:�(a)�300,�(b)�500�and�(c)�1000.��

  • � � �

    22�

    1000 800 600 400 200 0

    Si2pSi2s

    700�oC

    600�oC

    500�oC

    450�oC

    I3d

    O�KLL Ta4s Ta4p Ta4fO1s

    Ta4dC1s

    Cou

    nts�

    [a.u

    .]

    Binding�energy�[eV]630 620 610

    630.8619.4

    I3d�region

    700�oC

    600�oC

    500�oC

    450�oC

    Binding�Energy�[eV]

    No� iodine� could� be� detected� in� the� deposited� Ta2O5� films� by� XRFS.� XPS� was�used� to� confirm� these� results� and� also� to� check� the� O/Ta� ratio� and� the� chemical�bonding�state.�As�can�be�seen� in� the�XPS�spectra� (Figure�3.5)�no� iodine�could� be�detected�above� the�background�level� in�films�deposited�at�500�°C�and�higher.�Two�weak� I3d� peaks� could� be� observed� from� the� sample� deposited� at� 450� °C.� These�peaks� correspond� to� an� iodine� contamination� level� of�

  • � � �

    23�

    0.8 1.2 1.6 2.0 2.4 2.8 3.2 3.6 4.0

    a)

    b)

    Inte

    nsity

    �[a

    .u.]

    2θ�[o]

    Figure� 3.6.� XRR� curves� of� two� Ta2O5� films� on� Si(1�0�0)� substrates� with� different�thickness,�a)�12�nm�and�b)�20�nm.�Please�note�the�logarithmic�intensity�scale.��

    Some�typical�XRR�curves�are�shown�in�Figure�3.6.�As�can�be�seen�these�curves�exhibit�clear�thickness�fringes�due�to�constructive�and�destructive�interference�of�the�X-ray� beam� by� reflectance� from� a� uniform� surface� and� the� well-defined�film/substrate� interface.� The� XRR-curves� were� taken� by� using� a� line� focusing�geometry,� and� the� results� indicate� that� the� films� have� uniform� thickness� over� the�entire�substrate.�

    In�analogy�with�the�results�of�QCM�studies�of�TiO2�films�deposited�by�ALD�from�TiI4�and�O2��[19]�a�possible�growth�mechanism�could�be�described�as�follows:�

    TaI5(g)���TaIx(ads)�+�(5-x)/2�I2(g)� � (3.1)��

    TaIx(ads)�+�5/4�O2(g)���1/2�Ta2O5(s)�+�x/2�I2(g)� (3.2)��

    The� increasing� growth� rate� observed� at� higher� deposition� temperatures� might�then�be�explained�by� that�insufficient�energy� is�available�at� lower� temperatures�for�the� surface� reaction� (3.2)� to� occur� to� completion� during� the� O2� pulse� (t3),� i.e.,� the�process� is� thermally� activated.� � The� decrease� in� growth� rate� above� 625� °C� can� be�explained�by�an�increasing�rate�of�desorption�of�surface�species�(TaIx)�during�the�N2�purging� (t2).� Assuming� that� the� oxidation� of� TaIx(ads)� surface� species� at� low�temperatures� is� not� fully� realised� during� the� O2� pulse,� might� explain� why� small�amounts� of� iodine� was� detected� for� films� deposited� at� 450� °C� but� not� at� higher�deposition�temperatures.�

  • � � �

    24�

    Figure� 3.7.� Comparison� of� the� growth� rate� of� Ta2O5� on� Si(1�0�0)� substrates�employing� TaI5� and� two� different� oxygen� precursors� as� a� function� of� deposition�temperature�(middle).�SEM�pictures�of�Ta2O5�films�deposited�by�using�H2O/H2O2�at�250�°C� (left)� and� O2� at� 600�°C� (right)� as� oxygen� precursors.� In� both� cases� 1000�deposition�cycles�were�applied�on�Si(1�0�0)�substrates�

    It� is� interesting� to� compare� the� results� of� the� present� study� with� the� earlier� ALD�studies� using� the� TaI5-H2O/H2O2� precursor� combination� [42,� 70].� In� these� studies�Ta2O5�films�were�deposited�in�the�temperature�region�of�250�to�400�°C�on�Si(1�0�0)�substrates,�and�also�here�no�residual�iodine�could�be�detected�in�the�deposited�films.�The�films�were�found�to�be�amorphous�at�low�temperatures�and�partly�crystallised�at�350�°C.� The� growth� rate� decreased� from� 0.9�Å/cycle� (250�°C)� to� 0.6�Å/cycle�(400�°C)� in� the� investigated� temperature�range� (Figure�3.7,�middle).�These�growth�rates�are�thus�considerably�smaller�than�those�observed�in�the�TaI5-O2�process�(up�to�1.7�Å/cycle),�but�it�should�be�observed�that�the�values�here�were�obtained�at�higher�temperatures.� In� the� case� of� the� TaI5-H2O/H2O2� process� a� further� increase� in�temperature� will� most� certainly� result� in� an� increased� desorption� of� surface� -OH�groups�resulting� in�a�decreased�growth� rate.�Furthermore,� the�polycrystalline� films�of� the�present�study�have�a� larger�surface�area�compared�to� the�smooth�amorphous�films� of� the� earlier� study.� This� will� result� in� more� available� adsorption� sites,� and�accordingly�in�a�higher�growth�rate.��

    The� difference� in� surface� morphology� was� clearly� seen� by� SEM.� Figure� 3.7�shows�examples�of� the�different�surface�morphologies�from�films�deposited�by� the�two�processes�at�different�temperatures,�250�°C�(left)�resp.�600�°C�(right),�but�using�the�same�number�of�deposition�cycles�(1000).�The�H2O/H2O2�process�gives�smooth�amorphous� films� at� 250�°C� and� it� is� hard� to� see� any� surface� features� at� 100�k�magnification.� However,� in� the� case� of� the� O2� process� a� much� rougher� surface�morphology�is�revealed�since�the�films�are�polycrystalline.�

    250 300 350 400 450 500 550 600 650 700

    0.00.20.40.60.81.01.21.41.61.82.02.22.4

    CrystallineAmorphous

    TaI5-O

    2TaI5-H

    2O/H

    2O

    2

    Gro

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    �R

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    [Å/C

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    Deposition�Temperature�[oC]200�nm� 500�nm�

  • � � �

    25�

    3.2� �HfO2�In�ALD�of� HfO2� the� following�carbon� free�precursors�have�been�employed:� HfCl4�[71-73],� HfI4� [44]� and� Hf(NO3)4� [74].� The� oxygen� source� has� been� H2O� with� the�exception�of�the�HfI4�process,�where�a�mixture�of�H2O�and�H2O2�was�employed.�In�one�of� the�studies�of� the�chloride�based�process� the�growth�rate�was�0.5�Å/cycle�at�500� °C� and� the� films�did�not�have� any� chlorine� residues� (

  • � � �

    26�

    400 450 500 550 600 650 700 7500.0

    0.2

    0.4

    0.6

    0.8

    1.0

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    1.4 �pO

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    �=�13�Pa�p

    O2

    �=�40�Pa

    Gro

    wth

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    [nm

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    Growth�Temperature�[oC]

    0 200 400 600 800 10000

    200

    400

    600

    800

    1000

    1200

    Thi

    ckne

    ss�[Å

    ]

    Number�of�Cycles

    �pO

    2

    �=�13�Pa�p

    O2

    �=�40�Pa

    Figure� 3.8.� Growth� rate� as� a� function� of� growth� temperature� for� two� different� O2�partial�pressures,�13�and�40�Pa.��

    Two�thickness�series�were�deposited�at�620�°C,�one�using�the�lower�(13�Pa)�and�the�other�using�the�higher�(40�Pa)�O2�partial�pressure�(Figure�3.9).�The�growth�rate�was� found� to� be� proportional� to� the� number� of� deposition� cycles� in� both� cases.�Furthermore,� it�can�be�seen� that� for� the�higher�O2�partial�pressure,� the�growth�rate�increased�by�approximately�70�%.��

    Figure� 3.9.� The� thickness� of� HfO2� films� deposited� at� 620�°C� by� applying�different�number�of�deposition�cycles�and�at�two�different�O2�partial�pressures,�13�and�40�Pa.�The�dashed�lines�are�linear�fits�as�a�guide�for�the�eye.��

  • � � �

    27�

    30 35 40 45 50 55

    pO

    2

    �=�13�Pa

    2θ�=�50.4o2θ�=�30.4o

    500�oC

    570�oC

    620�oC

    590�oC

    660�oC

    Inte

    nsity

    �[a

    .u.]

    2θ�[�o�]30 35 40 45 50 55

    pO

    2

    �=�40�Pa

    2θ�=�50.4o

    620�oC

    500�oC

    400�oC2θ�=�30.4o

    Inte

    nsity

    �[a

    .u.]

    2θ�[�o�]

    The� observed� increased� growth� rate� with� increasing� O2� partial� pressure� can� be�explained�as�follows.�With�the�chosen�pulse�time�(2�s),�the�lower�O2�partial�pressure�(13�Pa)�is�insufficient�for�O2�to�react�with�all�adsorbed�HfI4�and�a�low�growth�rate�is�accordingly� obtained.�In�contrast,�using� the�higher�O2�partial�pressure� (40�Pa),� the�surface� reaction� with� the� adsorbed� HfI4� is� completed� resulting� in� a�higher� growth�rate.�A�further�increase�in�O2�partial�pressure�will�not�affect�the�growth�rate,�since�all�HfI4�has�already�reacted�at�40�Pa.�It�was�actually�observed�that�26�Pa�was�enough�in�later�film�deposition�experiments.�

    3.2.2� Structure�and�composition�TOF-ERDA� revealed� no� iodine� or� hydrogen� incorporation� in� the� films� that� were�grown�at�500� to�750�°C�using�a�O2�partial�pressure�of�40� Pa.�This�means� that� the�concentration� of� those� impurities� did� not� exceed� 0.1� at.� %.� The� main� impurity�detected�in�the�films�was�zirconium.�The�zirconium�concentration�was�1.2-1.5�at.%.�However,� the�oxygen� to�metal�(Hf�+�Zr)�atomic�ratio�was�equal� to�2.0�±�0.2� in�all�films� studied� in� this� work.� It� is� worth� mentioning� that�no� iodine� was� detected� by�XRFS� in�any� of� the�deposited� HfO2� films.�Since� iodine�has�not�been� found� in� the�films�deposited�using�a�lower�O2�partial�pressure,�which�resulted�in�a�lower�growth�rate,�it�must�be�assumed�that�even�though�not�all�adsorbed�HfIx�reacts�with�O2�during�the�first�O2�pulse�it�does�react�during�the�subsequent�O2�pulses.��

    Figure� 3.10.� The� GI-XRD� patterns� of� HfO2� films� grown� at� different� substrate�temperatures� using� a� low� (left)� and� a� high� (right)� O2� partial� pressure� and� by�applying� 1000� deposition� cycles.� All� peaks� can� be� attributed� to� monoclinic� HfO2�except�those�indicated�by�arrows.�Please�note�the�logarithmic�intensity�scale.��

    According�to�GI-XRD�(Figure�3.10)�the�dominating�phase�in�all�depositions�was�monoclinic�HfO2.�Phase�pure�monoclinic�HfO2�could�be�grown�above�620�°C�using�a� low�O2�partial�pressure�and�above�500�°C�using�a�high�O2�partial�pressure�when�applying�1000�deposition�cycles.�For� lower�deposition� temperatures,�a�peak�which�

  • � � �

    28�

    did�not�belong� to� the�monoclinic�phase,�was�observed�at�2θ�=�30.4°.�At� this�angle,�the�tetragonal�and�cubic�modifications�of�HfO2�should�give�a�1�1�1�reflection,�which�is� the� strongest� one� for� non-textured� samples� of� both� modifications.� At� other�diffraction� angles,� the� possible� reflections� from� both� HfO2� modifications� give�overlap�with�the�monoclinic�reflections.��

    Even� though� phase� pure� monoclinic� HfO2� films� could� be� grown� at� 620� °C� by�applying� 1000� deposition� cycles� and�a�high� O2� partial� pressure,� phase�analysis� by�GI-XRD�for�thinner�films�grown�at�this�temperature�showed�a�clear�presence�of�the�peak�at�2θ�=�30.4°�for�both�low�and�high�O2�partial�pressures.�The�relative�amount�of�metastable�phase,�characterized�by�the�relative�intensity�of�the�peak�at�2θ�=�30.4°,�is�generally� higher� for� thinner� films� and� the� relative� intensity� of� this� peak� decreases�with� film� thickness.� This� might� be� explained� by� that� a� metastable� HfO2� phase� is�formed�during�the�initial�stages�of�the�film�growth.�Furthermore,�the�amount�of�this�phase� initially� formed� is�mainly�dependent�on�growth� temperature.�The�metastable�HfO2�phase�can� then�be� transformed� into�monoclinic�HfO2�as� the�growth�proceeds�and� the� number� of� deposition� cycles� employed� i.e.,� the� duration� of� the� growth�process,� governs� the� amount� of� metastable� phase� left� when� the� growth� is� ended.�Figure� 3.11� shows� the� integrated� peak� intensities� for� films� deposited� at� 620�°C�at�both� low� and�high� O2� partial� pressures� as� a� function� of� the� number� of� deposition�cycles,� which� is� proportional� to� the� film� thickness.� It� can�clearly� be� seen� that� the�abundance� of� the� metastable� phase� is� decreasing� as� more� deposition� cycles� are�applied�and� that� the� relative�amount�of�metastable�phase� is� lower� in� the�case� of�a�higher�O2�partial�pressure.���

    Figure�3.11.�The�peak� intensities�of� the�monoclinic�–1�1�1�peak�(2θ� �=�28.4°�)�and�the�peak�at�2θ�=�30.4°�for� films�deposited�with�different�number�of�cycles�and�two�different�O2�partial�pressures.��

    200 400 600 800 10000

    50

    100

    0

    10

    20

    �Integrated�P

    eak�Intensity

    �[a.u.]Inte

    grat

    ed�P

    eak�

    Inte

    nsity

    �[a

    .u.]

    No.�of�Cycles

    �-1�1�1�Monoclinic,�p(O2)�=�40�Pa

    �-1�1�1�Monoclinic,�p(O2)�=�13�Pa

    �2θ�=�30.4o,�p(O2)�=�40�Pa

    �2θ�=�30.4o,�p(O2)�=�13�Pa

  • � � �

    29�

    Since� it�was�obvious� that�the�metastable�phase�appeared� in�the�beginning�of� the�growth�process,�it�was�decided�to�monitor�the�relative�abundance�of�this�phase�as�a�function� of� film� thickness.� In� Figure� 3.12� the� relative� intensities� of� the� -1� 1� 1�monoclinic�peak�and� the�peak�at�2θ�=�30.4°� is�plotted�as�a�function�of� the� incident�angle�of�the�X-ray�beam.�The�penetration�depth�below�φc�can�be�calculated�according�to�the�following�expression�[75]:��

    t’�=�λ�/�[2π(φc2-φ2)1/2]� t’�=�penetration�depth�[Å]�� � λ�=�wavelength�[Å]�� � φc�=�critical�angle�for�total�reflection�[rad]�

    � � φ�

    =�incident�angle�[rad]��

    Figure�3.12.�The�intensities�of�the�-1�1�1�monoclinic�peak�and�the�peak�at�2θ�=�30.4°�is�plotted�as�a�function�of�the�incident�angle�of�the�X-ray�beam�for�a�film�deposited�by�applying�350�cycles�(420�Å)�at�600�°C.�The�solid�line�represents�the�penetration�depth�below�the�critical�angle�(αc(HfO2)�=�0.32°,�measured).�The�right�picture�is�a�schematic�drawing�of�the�phase�distribution�in�the�film.�����

    For� incident�angles�(φ)�below�φc� information� is�collected� from� the�surface� layer�of�the�film.�As�can�be�seen�in�Figure�3.12�the�metastable�peak�at�2θ�=�30.4°�gives�no�or� very� low� intensity� in� this� region� (φ� =� 0� -� 0.32°).� The� curve� that� represents� the�-1�1�1�monoclinic�peak�has�a�clear�maximum�around�φc,�which�means�that�this�phase�is�present�in�the�upper�region�of�the�film.�The�intensity�of�the�monoclinic�peak�does�not� decrease� to� a� large� extent� as� the� incident� angle� is� increased� above� φc� and� the�penetration�depth�of�the�x-rays�increases,�which�means�that�the�monoclinic�phase�is�present�throughout�the�film.�However,�the�peak�from�the�metastable�phase�does�not�show�a�maximum�in�intensity�for�incident�angles�around�φc�and�is�hence�not�present�

    0.0 0.2 0.4 0.6 0.8 1.0

    φ�[�o�]

    Inte

    grat

    ed�In

    tens

    ity�[a

    .u.]

    0

    100

    200

    300

    400

    500

    x�10

    2θ�=�30.4o

    -1�1�1�Monoclinic

    φc(HfO

    2)

    Pen

    etra

    tion�

    Dep

    th�[Å

    ]monoclinic�

    HfO2�.�.�.�.�

    monoclinic�+�metastable�

    HfO2�

    SiO2(20�Å)/Si(100)sub�

  • � � �

    30�

    a� b� c�

    in� large� amounts� in� the� top� layer� of� the� film.� This� study� thereby� supports� the�previous�statement�that�the�metastable�forms�initially�during�film�growth.�

    When� comparing� the� XRD� patterns� from� films� of� similar� thickness� a� peak�broadening�of�the�monoclinic�reflections�was�observed�for�films�with�higher�relative�amount�of�metastable�phase,�which�is�an�indication�of�relatively�smaller�grain�sizes.�This�observation�might�be�explained�by�that�crystallite�sizes�affect�the�phase�stability�i.e.,� that� the� metastable� phase� is� stabilised� by� smaller� grain� sizes.� Since� it� is�generally� known� that� the� grain� size� increases� with� temperature� this� would� explain�why� smaller� amounts� of� the� metastable� phase� are� formed� when� the� growth�temperature� is� increased.� There� is�no� obvious� difference� in� relative� peak� intensity�when�comparing�films�with�similar�thickness�but�grown�by�using�different�O2�partial�pressures.�Therefore�the�O2�partial�pressure�primarily�affects�the�growth�rate�for�the�growth�conditions�investigated.���

    Figure�3.13.�Transmission�electron�images�of�the�HfO2�films�grown�on�HF-stripped�silicon�at�a)�500�°C,�b)�620�°C�and�750�°C�with�an�oxygen�partial�pressure�of�40�Pa.�In�all�three�pictures�the�layer�order�is�HfO2/SiO2/Si(1�0�0)�going�from�left�to�right.��

    A�series�of�samples�was� investigated�by�TEM�(Figure�3.13)� to�obtain� images�of�the�cross-section�between�the�Si(1�0�0)�substrate�and�the�HfO2�film.�As�can�be�seen�in� the� TEM� pictures� there� is� a� 20�Å� thick� interface� layer� and� the� thickness� is�independent�of�the�deposition�temperature�employed�(500�to�750�°C)�when�using�an�O2�pressure�of�40�Pa.�It�is�clear�that�the�interface�layer�consists�of�SiO2�since�Hf�was�not�detected�in�the�EDX�spectra�taken�in�the�vicinity�of�the�Si�surface�while�Si�and�O�was�detected�in�this�region.�

    3.2.3� Electrical�characterisation�Some� of� the� electrical� data� obtained� from� Al/HfO2/Si� structures� are� listed� in�Table�3.1.�The�leakage�current�for�these�structures�remained�quite�low,�in�the�order�of�nanoamperes�until�dielectric�breakdown.�The�breakdown�field�decreases�with�an�increase�in�growth�temperature.�It�is�obvious�that�films�deposited�using�a�higher�O2�

  • � � �

    31�

    -4 -3 -2 -1 0 1 2 3 4 50.0

    0.1

    0.2

    0.3

    0.4

    f�=�100�kHzpO2

    �=�0.1�Torr

    HfI4-O2

    Al/HfO 2/p-Si(100)

    �570�oC,�65�nm

    �620�oC,�79�nm

    �660�oC,�66�nm

    �755�oC,�167�nm

    Ca

    pa

    cita

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    ,�n

    F

    Bias�voltage,�V

    -4 -3 -2 -1 0 1 20.0

    0.1

    0.2

    0.3

    0.4

    0.5

    0.6

    0.7

    f�=�100�kHz

    111�nm700�oC

    54 �nm750 �oC

    26 �nm620�oC

    pO 2�=�0.3�TorrHfI4-O 2

    Al/HfO2/p-Si(100)

    Ca

    pa

    cita

    nce

    ,�n

    F

    Bias�voltage,�V

    partial� pressure� are� more� leaky� and� less� resistant� to� breakdown� than� the� films�deposited�using�the�lower�O2�partial�pressure�when�comparing�the�field�required�for�dielectric�breakdown.�

    Table�3.1.�Electrical�data�of�selected�HfO2�films.�

    p(O2)��

    Tgrowth��

    Breakdown�field��

    Effective�permittivity��

    Thickness�[Pa]� [°C]� [MV/cm]� � [Å]�13� 570� 2.1� 13.8� 540�13� 620� 1.6� 12.1� 790�13� 660� 1.0� 10.2� 660�13� 700� 0.7� 14.0� 1000�40� 620� 1.5� 15.5� 260�40� 700� 0.5� 13.0� 1110�40� 750� 0.6� 13.0� 540�

    Figure� 3.14.� Capacitance-voltage� curves� measured� on� Al/HfO2/Si� capacitators� at�100�kHz,�where�the�O2�partial�pressure�was�a)�13�Pa�and�b)�40�Pa.��

    a)�

    b)�

  • � � �

    32�

    Figure�3.14�a�demonstrates�selected�capacitance-voltage�curves�of�the�Al/HfO2/p-Si(1�0�0)�structures,�where�the�HfO2�was�grown�using�an�O2�pressure�of�13�Pa.�It�can�be�seen� that� the�curves�demonstrate�counter�clockwise�hysteresis�and� the� flat-band�voltage� is� shifted� towards� positive� bias� voltages� upon� increase� in� HfO2� growth�temperature.�It�is�to�be�noted�that�the�flat-band�voltage�in�the�case�of�high�interface�quality� should� occur� at� -1� V� for� the� Al� and� p-Si� electrodes.� The� existence� of� the�hysteresis� refers� to� rechargeable� traps� in� the� oxide� layer,� with� the� charge� carriers�being�mainly�electrons�that�become�ejected�from�the�oxide�at�strongly�negative�bias�voltages.�[76]�The�shift�in�the�flat�band�voltage�indicates�that�the�amount�of�negative�charge�in�the�oxide�increases�with�the�processing�temperature.�The�flat-band�voltage�shift�is�also�accompanied�by�the�stronger�stretchout�of�the�C-V�curves.��

    In� the� case� of� the� HfO2� grown� using� an� O2� dose� of� 40� Pa,� the� C-V� curves�demonstrate�very�narrow�counter�clockwise�hysteresis�or�no�hysteresis�at�all,�and�the�flat-band� voltage� is� shifted� more� strongly� towards� positive� bias� voltages� at� the�lowest�growth� temperatures�(Figure�3.14�b).�Upon� increase� in�the�HfO2�processing�temperature,�the�flat-band�voltage�shifts�to�–1V,�approximately,�which�is�consistent�with� the� theoretically� assumed� position� [77].� Thus,� in� the� case� of� higher� oxygen�pressure,� the� flat-band� voltage� behaved� in� an� opposite� way� when� the� deposition�temperature� is� changed� compared� to� the� temperature� series� of� samples� obtained�using�the�low�oxygen�pressures.�

    The� observations� reported� above� might� be� connected� to� crystallite� sizes� at� the�interface.� The� metastable� phase� is� as� mentioned� (section� 3.2.2.)� stabilised� by�relatively�smaller�crystallite�sizes�compared�to�the�monoclinic�phase.�Large�amounts�of�metastable�phase�may�form�at�the�interface�initially�during�growth,�depending�on�the� growth� temperature.� Depending� on� how� thick� the� film� is� allowed� to� grow� a�certain�amount�of�metastable�phase�with�relatively�smaller�crystallite�size�remains�at�the�interface.�

    The�electrical�characterisations�that�has�been�presented�suggests�that�films�grown�at�lower�partial�O2�pressures�and�at�lower�growth�temperatures�has�a�higher�ability�to�withstand�dielectric�breakdown.�One�reason�for�this�observation�might�be�connected�to� the� fact� that� these� films� have� a� higher� relative� amount� of� relatively� smaller�crystallites� at� the� film/substrate� interface� as� seen� by� the� GI-XRD� investigation.�Furthermore,� these� films� also� possess� a� wider� hysteresis� in� the� CV-curves,� which�might� be� connected� to� charge� trapping� in� the� film/substrate� interface,� where� the�smaller�crystallites�are�present�(section�3.2.2).�

  • � � �

    33�

    Si(100)�

    100��

    SiO2�

    Figure� 3.15� Transmission� electron� image� of� a� HfO2� films� grown� by� applying� 50�deposition�cycles�on�HF-stripped�silicon�at�620�°C�with�an�oxygen�partial�pressure�of�40�Pa.�

    Studying�the�films�deposited�at�620�°C�and�above�by�using�the�higher�O2�partial�pressure,�which�have�no�detectable�amounts�of� the�metastable�phase�(Figure�3.10),�can� build� additional� support� to� these� arguments.� Their� CV-curves� (Figure� 3.14)�possesses�no�or�minute�signs�of�hysteresis�and�they�are�relatively�easier�affected�by�dielectric�breakdown�(Table�3.1).�According�to� the�GI-XRD� investigation� the�grain�size�of�the�monoclinic�phase�is�typically�about�300�Å�in�thicker�films�(according�to�Scherrers� formula).� Figure� 3.15� shows� a� TEM� micrograph� of� the� film/substrate�interface�for�a�film�deposited�by�applying�50�deposition�cycles�at�620�°C.�It�can�be�seen�that�the�SiO2�layer�has�been�formed�already�at�this�stage�of�growth�and�that�the�crystallite�size,� judging�by� the� three�visible�crystallites,� is�approximately�100�Å.�If�the�growth�had�been�allowed�to�continue�the�crystallites�would�probably�increase�in�size�as�for�the�thic