Failure and Fatigue of Nanostructured Metals | Fracture | Deformation (Engineering)

Good review paper on failure and fatigue of nanostructured metals
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  By invitation only: overview article Failure of metals III: Fracture and fatigue of nanostructured metallicmaterials Andr  e Pineau  a ,  * , A. Amine Benzerga  b ,  c , Thomas Pardoen  d ,  * a Mines ParisTech, Centre des Mat   eriaux, CNRS UMR 7633, BP 87, 91003 Evry, France b Department of Aerospace Engineering, Texas A & M University, College Station, TX 77843, USA c Department of Materials Science and Engineering, Texas A & M University, College Station, TX 77843, USA d Institute of Mechanics, Materials and Civil Engineering, Universit   e catholique de Louvain, B-1348 Louvain-la-Neuve, Belgium a r t i c l e i n f o  Article history: Received 20 March 2015Received in revised form18 July 2015Accepted 18 July 2015Available online 11 August 2015 Keywords: DuctilityFractureFatigueMetalsNanocrystallineThin  󿬁 lmsNeckingDamage a b s t r a c t Pushing the internal or external dimensions of metallic alloys down to the nanometer scale gives rise tostrong materials, though most often at the expense of a low ductility and a low resistance to cracking,with negative impact on the transfer to engineering applications. These characteristics are observed,with some exceptions, in bulk ultra- 󿬁 ne grained and nanocrystalline metals, nano-twinned metals, thinmetallic coatings on substrates and freestanding thin metallic  󿬁 lms and nanowires. This overview en-compasses all these systems to reveal commonalities in the srcins of the lack of ductility and fractureresistance, in factors governing fatigue resistance, and in ways to improve properties. After surveying thevarious processing methods and key deformation mechanisms, we systematically address the currentstate of the art in terms of plastic localization, damage, static and fatigue cracking, for three classes of systems: (1) bulk ultra- 󿬁 ne grained and nanocrystalline metals, (2) thin metallic  󿬁 lms on substrates, and(3) 1D and 2D freestanding micro and nanoscale systems. In doing so, we aim to favour cross-fertilizationbetween progress made in the  󿬁 elds of mechanics of thin  󿬁 lms, nanomechanics, fundamental researchesin bulk nanocrystalline metals and metallurgy to impart enhanced resistance to fracture and fatigue inhigh-strength nanostructured systems. This involves exploiting intrinsic mechanisms, e.g. to enhancehardening and rate-sensitivity so as to delay necking, or improve grain-boundary cohesion to resistintergranular cracks or voids. Extrinsic methods can also be utilized such as by hybridizing the metalwith another material to delocalize the deformation - as practiced in stretchable electronics. Fatiguecrack initiation is in principle improved by a  󿬁 ne structure, but at the expense of larger fatigue crackgrowth rates. Extrinsic toughening through hybridization allows arresting or bridging cracks. The contentand discussions are based on experimental, theoretical and simulation results from the recent literature,and focus is laid on linking microstructure and physical mechanisms to the overall mechanical behavior. ©  2015 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. 1. Introduction While the companion overviews [1,2] deal with recent researchadvances in the fracture of industrial metallic alloys involvingtypicalmicrostructuredimensionswithinthetensofmicronsrangeas used in traditional macroscopic components, the present over-view addresses the state of the art and future perspectivesregarding the physics and mechanics of fracture in metallic mate-rials with nano-scale internal or external dimensions. Severalreviews have been dealing in recent years with the mechanicalbehavior of these systems, e.g. [3 e 8], with a focus on the defor-mation mechanisms controlling the extreme strengths that can beattained,butoftenwithlimitedattentiontofracture.Theresistanceto fracture is usually reported tobeweak, as quanti 󿬁 ed bya limitedductilityand/orlowfracturetoughness.Thisisakeyissuethatmustbe solved in order to open the range of possible viable applicationsof nanostructured metallic materials. Fig. 1 describes the coverageof the paper; it deals with bulk  “ 3D ”  macroscopic metals involvingsubmicroncharacteristicinternalmicrostructurefeaturesaswellasmetallic structures involving at least one submicron dimension, i.e.both  “ 2D ”  󿬁 lms or coatings and  “ 1D ”  wires, rods and beams whichinvolve submicron characteristic internal microstructure sizes aswell. The fracture of 1D and 2D structures is more complex in a *  Corresponding authors. E-mail addresses: (A. Pineau), (T. Pardoen). Contents lists available at ScienceDirect Acta Materialia journal homepage: ©  2015 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Acta Materialia 107 (2016) 508 e 544  sense, as both the microstructural and component characteristiclengths generally affect it. Incidentally, covering all three types of material systems enableslinkstobemadebetweenthecommunitythat develops bulk nanostructured alloys for potential structuralapplications and that working on nano- and thin- 󿬁 lm mechanicsfor electronics, coatings and nanotechnology. In order to limit thescope, only room temperature static fracture and fatigue of crys-talline metals will be addressed with no consideration for otherimportant issues related to environment, high or low temperatureeffects or to metallic glasses. Fracture is addressedfromthe generalperspective of failure analysis in the presence or not of a pre-crack,under static or cyclic conditions. The ductility is set either by theonset of plastic localization or by the onset of cracking. Fatigue andfracturearetreatedintermsof theresistancetocrackinitiationandpropagation from a pre-existing crack.Before entering into the topic of fracture, Section 2 sets a seriesof basic elements on processing and deformation mechanisms. Theprocessing methods are brie 󿬂 y covered while presenting the spe-ci 󿬁 c materials and systems addressed in this overview. The secondpart of Section 2 provides a survey of the mechanisms that controlthe strength in metals with nanoscale dimensions. These aspectsare needed and must be understood as they dictate the ensuingmechanisms of localization, damage and fracture. Indeed, theessence of the  “ local approach ”  to fracture that traverses the threeoverviews is to address failure mechanisms in their connectionwith the microstructure and the physical phenomena. Damage andcracking are the results of the speci 󿬁 c deformation mechanismsdeveloping in this class of materials.The fracture and fatigue behavior of bulk 3D nanostructuredmetals is treated in Section 3. The research in this  󿬁 eld has beenintense over the last two decades, opening applications to a varietyof engineering  󿬁 elds where a high strength is the main perfor-mance of interest. One of the key issues with these materials, inaddition to the dif  󿬁 culty to transfer the processing routes toindustrialstructuralapplications,istheusuallylowductilitybothinterms of uniform elongation and fracture strain which negativelyimpact the forming capabilities and resistance to cracking whenusedinstructuralcomponents.Thelowstrainhardeningcapacityisthemainreasonforthelimitedresistancetoplasticlocalizationandfor the low uniform elongation. Regarding damage and cracking,the key issue is the predominance of intergranular decohesionmechanisms because of the extreme stress build-up on grain-boundaries (GBs) and the presence of processing defects. Work-arounds recently found to recover ductility by introducing het-erogeneities atmultiplescalesandbyfavouringrate sensitivity willbe analysed, for instance the nano-twinned (NT) metals [9] orhybrid nano-structuring strategies involving bimodal or multi-modal grain size distributions, multi-phase alloys, graded single ormultiphase systems, multi-layered metallic composites, generationofcontrolledinternal stressdistributions,andcombinationof thesesolutions.The fracture of 2D and 1D materials is addressed in Sections 4and 5. Fracture in thin metallic layers on substrate has probablybeen more studied than in bulk nanostructured metals, partlybecause these systems have been available for a long time owing towell-controlled deposition/coating methods. A seminal example isfound in galvanization layers (although usually not exhibitingsubmicron thickness). The mechanics of fracture of coatings onsubstrates has been rigorously established more than 20 years ago,as reviewed in [10]. Thin metallic layers are used in awide range of applications involving microelectronics, protective and/or func-tional coatings, and membranes. The ductility and fracture resis-tance of the  󿬁 lms is often a key property needed in applications. In 󿬂 exible electronic devices, stretchability is essential for preservingelectrical conductivity even under large distortions. In thin func-tional coatings, the ductility of the thin layer or multilayer must belarge enough to sustain forming operations performed on thesubstrate or to resist scratching or impact conditions, as well as Fig.1.  Description of the different UFG and nanocrystalline metallic materials, single phase or hybrids, addressed in the paper from the viewpoint or their failure behavior, involvingbulk systems (Section 3), thin  󿬁 lms and coatings on substrate (Section 4) and freestanding thin  󿬁 lm and nanowires (Section 5).  A. Pineau et al. / Acta Materialia 107 (2016) 508 e 544  509  thermal loadings. Micro- or nano-electromechanical systems(MEMS/NEMS) devices and other microsystems sometimes involvethin metallic  󿬁 lms either freestanding or lying on a substrate, andwhich can undergo signi 󿬁 cant straining under operation ormanufacturing. In microelectronic devices, large strains in thininterconnects result from thermal cycling and associated internalstress evolution. In all these applications, permanent plasticdeformation is not necessarily prohibited, but should developwithout or with limited micro-cracking to preserve the expectedfunction. The same problem as for bulk nanostructured systems of Section3emerges:mostmetallic 󿬁 lmsexhibitlimitedductility[11],even though recent counterexamples have been described in theliterature with thin freestanding metallic  󿬁 lms deforming by morethan 10%, e.g. [12,13]. For the sake of clarity, the fracture of metallic coatings while lying on a substrate is addressed in Section 4, andthefractureoffreestanding 󿬁 lmsorwiresinSection5.Eventhoughthe intrinsic failure mechanisms are usually similar, the experi-mental challenges and failure scenarios can be quite different. Inparticular, the  󿬂 exible electronics community has developed a se-ries of innovative extrinsic ductilization strategies, involving simi-larities with the hybrid approaches applied to bulk nanostructuredmetals. In the case of freestanding 󿬁 lms, we review recent progressabout in situ characterization offailure mechanisms witharealmof new observations and improved understanding.InthesethreecoreSections3 e 5,wewilladdressfailureintermsofthepropertiesthatareusedtoquantifytheresistancetofracture:(1) the  uniform elongation , in terms of the engineering strain  A  g  (%) or true strain  ε u  (/), which is the strain at the onset of diffuse necking;(2) the  true fracture strain  ε  f   (/) which is the effective plasticstrain attained in the broken region as measured from thearea reduction of section;(3) the total elongation  ε engf   (/)whichis extractedfromauniaxialtensile test as the total displacement at fracture divided byinitial length. This quantity is often reported in the literaturebutcanlead toconfusionas itisgaugelengthdependent andvery much dominated by the magnitude of   A  g  ;(4) the  true fracture stress  s  f   (MPa) which may be more mean-ingful than the true fracture strain when the behavior isbrittle or quasi-brittle;(5) the  fracture toughness  expressed by the critical stress in-tensity factor  K  Ic   (MPa   m), critical energy release rate  G Ic   (J/m 2 ), critical value of the  J   integral  J  Ic   (J/m 2 ), or the criticalCrack Tip Opening Displacement (CTOD)  d Ic   ( m m) (limitedhere to mode I crack opening conditions) which requires, tobe valid, samples with a sharp pre-crack and suf  󿬁 cientlylarge dimensions to represent an intrinsic, geometry-independent material fracture resistance index. Thesevalues are usually de 󿬁 ned at cracking initiation except if indicated otherwise. The meaning and relevance of theseconcepts in the context of submicron components will becritically discussed;(6) the  resistance to fatigue  expressed by the  endurance limit   s end (MPa) which is the stress amplitude above which fracture byfatigue does not occur below a given high number of cycles(usually 10 6  10 7 cycles), or by the  D a / D N   ( m m) versus  D K  (MPa   m)responsewhichgivestheincrementofcrackgrowth D a  per loading cycle  D N   under a given amplitude of stressintensity factor  D K  .Even though a material is not physically broken when plasticlocalization occurs (as quanti 󿬁 ed with  A  g   or  ε u ), the integrity of thestructure is compromised and, from an engineering viewpoint, onemay consider that failure has been attained. The prediction of theintegrity of a structural component or the formability of a materialpart must be based on a series of failure criteria involving bothfracture from preexisting defect or not and plastic localization (andbuckling, though not discussed here). We will discuss how theseproperties play out in the systems described above, how theyconnectwiththemicrostructureandwiththevariouscharacteristiclengths involved. On the one hand, the small size makes the testingand characterization more complicated for some aspects, such asfor quantitative load measurement, manipulation and alignment,but,ontheotherhand,directobservationwithTEMonthefullscalespecimen can be performed and volumes similar to those analysedexperimentally can be treated with atomistic simulations ordislocation dynamics codes. For this, relevant recent literature onexperimental, computational and theoretical researches will bereviewed involving open questions. Even though the literature onfracture and fatigue of these systems is much less extensive thanthe literature dealing with deformation and strength aspects, it isstill wide enough and we do not claim a comprehensive treatment.Instead, we propose some possible lines of articulation for theunderlying questions and recent  󿬁 ndings. 2. Basics on processing, microstructures and deformationmechanisms  2.1. Processing of nanostructured metals This sub-section involves information both on the processingmethods and main characteristics of the micro- and nano-structures, with an emphasis on the elements that are needed tounderstand the different topics addressed in the paper. The termnano-crystalline (NC) will be reserved for grain sizes below~100 nm and ultra- 󿬁 ne grained (UFG) for grain sizes between~100 nm and ~1  m m. Classical metallic polycrystals will be referredas coarse grained (CG). A  󿬁 lm is considered as  “ thin ”  when thethicknessisbelow1  m mandananowireisde 󿬁 nedashavingamoreor less equiaxed section with a characteristic dimension below1  m m.  2.1.1. Bulk ultra  󿬁 ne grained and nanocrystalline metals The most common method to produce UFG and NC metallicalloys is through  Severe Plastic Deformation  (SPD), under high hy-drostatic pressure to avoid failure problems, sometimes at highstrain rates. Many reviews have been dedicated to this topic, e.g.[14 e 16]. There is a large number of possible routes to apply ultradeformation involving [16] equal-channel angular pressing (ECAP),high pressure torsion (HPT), accumulative roll bonding (ARB),multi-axial forging, twist extrusion, swaging and all sorts of de-rivatives of these processes. Depending on the details of the pro-cess,thegrainsizecanbere 󿬁 nedtoameandimensionbelow1  m m,and sometimes below 100 nm by a mechanism of grain fragmen-tationthroughdislocationcellsformation,seeanexampleinFig.2aofavery 󿬁 negrainedCumicrostructure obtainedbyECAP [17].SPDmethods have been successfully applied on most pure and alloyedmetals.In termsof hybrid nanostructured metals, if the initial grain sizeobtained by SPD is suf  󿬁 ciently small as in Fig. 2a, a bimodalstructure can be produced by heat treatments with the size of thelarge grains in the micron range, see the example of Cu in Fig. 2b[17]. The ARB process has been successfully applied to producelaminates of alternating Cu and Nb layers owing to repeated rollbonding, cutting, stacking and further rolling, e.g. [18,19]. The layerthicknesses range from ~10 to ~1000 nmwith up to meter scale in-plane dimensions. A variant of the ARB process named  “ repeatedpress and rolling ”  has been used to process nano-layered com-posites with alternating Cu/Fe, Cu/Ag, Fe/Ag and Cu/Nb  A. Pineau et al. / Acta Materialia 107 (2016) 508 e 544 510  compositions, e.g. [20 e 22]. Accumulative  “ extrusion, drawing andbundling ”  can generate nano- 󿬁 laments of Nb within a channeling,highly textured Cu matrix with a controlled hierarchical structuresee e.g. [23]. Shot peening can be also considered as a SPD processas it heavily deforms under high strain rates the near surface re-gions of metals, leading to a hybrid structure with very  󿬁 ne grainsizes at the surface and a gradient towards to core with largergrains [24,25].Recently, new routes involving  advanced thermo-mechanicaltreatments  have been proposed to produce nanostructuredmetals, especially steels, with processing routes compatible withindustrial practice. In particular, there is a growing interest in thedevelopment of new advanced high strength steels (AHSS) withenhanced combination of strength and ductility. UFG dual phase(UFG DP) steels can be produced using two types of concept: theQ  & P and the inter-critical heat treatments. In both cases, meta-stable retained austenite (RA) can transform into martensites  ε (hcp) and  a /or  a 0  (bcc or BCT) leading to a transformation inducedplasticity (TRIP) effect. The RA phase may also remain stable andgive rise to mechanical twinning due to its low stacking fault en-ergy (SFE), especially in high Mn (~10%wt) steel. The UFG steels arethus complex hybrid multiscale systems. The UFG DP steels areknown for a long time (see e.g. Jin et al. [26]). Their microstructuremayalsobemetininter-criticallyheat-affectedzonesofmulti-pass Fig. 2.  Set of UFG, NC and thin  󿬁 lm metallic microstructures; (a) NC Cu produced by ECAP [17]; (b) bimodal structure produced after heat treatment from the microstructure shownin (a); (c) TEM micrographs of the submicron thick  󿬁 lm type retained austenite in BAT (see de 󿬁 nition on the text) from [29]; (d) High Mn steel with very  󿬁 ne-grained austenite-ferrite ( < 400 nm) duplex microstructure obtained bycombining rolling and intercritical annealing, from [30]; (e) nanotwinned Cu produced byelectrodeposition, from [42]; (f) TEM bright  󿬁 eld micrograph of a FIB cross section of a 300 nm thick Pd  󿬁 lm with 30 nm elongated columnar grains exhibiting a  〈 111 〉  texture parallel to the normal of the  󿬁 lm [51]; (g)TEM bright  󿬁 eld micrograph of a 40 nm Cu/40 nm Nb multilayer showing {111}Cu||{110}Nb  󿬁 ber texture [228].  A. Pineau et al. / Acta Materialia 107 (2016) 508 e 544  511
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