A review of high-strength nanolaminates and evaluation of their properties

doi:10.1016/j.jmst.2020.03.011

Abstract

Abstract:

Nanolaminates are composed of nanoscale-thick alternating layers of different materials and their properties are dependent on the individual layers, the layer thickness and the interfaces between the layers. Nanolaminates composed of cubic crystal structured metals are usually ductile compared to nanolaminates containing hexagonal crystal structured metals. Mechanical properties such as strength and hardness of nanolaminates increase with a decrease in individual layer thickness down to a few nanometers and they become independent when the thickness of individual layers is less than a couple of nanometers. This review provides a detailed analysis of the effects of individual layer thickness and the interface structures on the strength and the strengthening mechanisms of nanolaminates, their ductility and fracture behavior in terms of structural variations including grain morphologies, nanotwins, amorphous phases and crystal structures of the layers. The principles for designing nanolaminates with exceptionally high mechanical and physical properties and their fabrication are also highlighted. Some contradictory issues such as strengthening mechanisms, elastic modulus dependency on individual layer thickness and the effect of a thin amorphous layer on the strength are discussed. This review also provides future research directions in designing the high-strength nanolaminates that will facilitate practical engineering applications through analyzing up-to-date research efforts.

Key words: Nanolaminates, Mechanical properties, Physical properties, Deformation mechanisms

Mohammad Nasim, Yuncang Li, Ming Wen, Cuie Wen. A review of high-strength nanolaminates and evaluation of their properties[J]. J. Mater. Sci. Technol., 2020, 50: 215-244.

Figures/Tables 39

Fig. 1. Schematic of physical vapor deposition using magnetron sputtering [11].

Fig. 2. Schematic of cyclic atomic layer deposition process.

Fig. 3. Schematic representation of: (a) single bath ED system and (b) dual-bath ED system.

Fig. 4. Schematic of CVD process.

Fig. 5. (a) SEM cross-section in backscattered mode illustrating multilayer structure of Al2O3/ZrO2 nanolaminate (dark layers correspond to Al2O3 and bright layers correspond to ZrO2); and (b) TEM lattice-fringe image showing interface between an alumina and a zirconia grain (zirconia and alumina are indicated with no intergranular phase) [57].

Fig. 6. (a) Schematic representation of ARB process and (b) typical bright-field TEM image of as-rolled ARB multilayered Cu/Nb nanocomposite [19].

Table 1 Summary of the advantages and limitations among different fabrication techniques of nanolaminates.

Nanolaminate fabrication techniques	Advantages	Limitations
Physical vapor deposition (PVD)	- better control of the layer composition and thickness during nanolaminate fabrication - magnetron sputtering can produce high-quality and low-impurity metal and ceramic nanolaminates - higher deposition rate - good adhesion between individual layers in nanolaminate structures - magnetron sputtering can operate at lower deposition temperatures even at room temperature, which overcomes the difficulties of using lower melting temperature substrates	- it requires high vacuum to fabricate nanolaminates
Atomic layer deposition (ALD)	- it can control the thickness on the atomic scale during nanolaminate fabrication - has option to select wide range of precursor materials - good for relatively high-density material deposition with low impurity	- it requires modest temperature (<350 °C) to fabricate nanolaminates. Outside of the temperature range poor deposition is occurred due to either lower deposition rate or thermal decomposition of precursors.
Electrodeposition (ED)	- technologically simple structure of the system - compositional and structural changes of nanolaminates can be made by varying the deposition parameters - dual-bath ED technique overcomes the limitation of choosing wide range alloy compositions	- deposition parameters can be varied during the deposition and adversely affect the nanolaminate structures - single-bath ED technique suffers due to the selection of limited materials - dual-bath ED technique arises complexity and undesirable surface reactions during transferring the substrate between two electrolytes - possibility of high impurity and blunt interface structures in nanolaminates
Chemical vapor deposition (CVD)	- PE-CVD is advantageous due to its low temperature deposition, even it can operate at ambient temperature	- limited usages of low melting temperature substrates in thermal CVD process - interfaces between the layers of nanolaminates are not sharp enough due to the diffusion at high temperatures - lower deposition rate than PVD
Accumulative roll bonding (ARB)	- favorable for continuous production of nanolaminates - low cost and high production rate	- atomically smooth and few nm alternate layers in nanolaminates may not be achieved

Table 1 Summary of the advantages and limitations among different fabrication techniques of nanolaminates.

Nanolaminate fabrication techniques	Advantages	Limitations
Physical vapor deposition (PVD)	- better control of the layer composition and thickness during nanolaminate fabrication - magnetron sputtering can produce high-quality and low-impurity metal and ceramic nanolaminates - higher deposition rate - good adhesion between individual layers in nanolaminate structures - magnetron sputtering can operate at lower deposition temperatures even at room temperature, which overcomes the difficulties of using lower melting temperature substrates	- it requires high vacuum to fabricate nanolaminates
Atomic layer deposition (ALD)	- it can control the thickness on the atomic scale during nanolaminate fabrication - has option to select wide range of precursor materials - good for relatively high-density material deposition with low impurity	- it requires modest temperature (<350 °C) to fabricate nanolaminates. Outside of the temperature range poor deposition is occurred due to either lower deposition rate or thermal decomposition of precursors.
Electrodeposition (ED)	- technologically simple structure of the system - compositional and structural changes of nanolaminates can be made by varying the deposition parameters - dual-bath ED technique overcomes the limitation of choosing wide range alloy compositions	- deposition parameters can be varied during the deposition and adversely affect the nanolaminate structures - single-bath ED technique suffers due to the selection of limited materials - dual-bath ED technique arises complexity and undesirable surface reactions during transferring the substrate between two electrolytes - possibility of high impurity and blunt interface structures in nanolaminates
Chemical vapor deposition (CVD)	- PE-CVD is advantageous due to its low temperature deposition, even it can operate at ambient temperature	- limited usages of low melting temperature substrates in thermal CVD process - interfaces between the layers of nanolaminates are not sharp enough due to the diffusion at high temperatures - lower deposition rate than PVD
Accumulative roll bonding (ARB)	- favorable for continuous production of nanolaminates - low cost and high production rate	- atomically smooth and few nm alternate layers in nanolaminates may not be achieved

Fig. 7. Schematic representation of nanolaminate grain morphologies of: (a) equiaxed grains, (b) compressed grains, and (c) superlattice columnar structures.

Fig. 8. HRTEM images of: (a) (111)fcc‖(110)bcc K-S interface in Cu/Nb nanolayered composite and (b) (111)fcc‖(110)bcc N-W interface in Cu/Nb nanolayered composite [88].

Fig. 9. (a) Typical microstructures of Ag/Cu nanolamellar composites processed by high-pressure torsion (HPT) (arrows indicate twins), (b) HRTEM images of profuse deformation of twinning and stacking faults (SFs) in Ag nanolayers, (c) a deformation twin nucleated from an interface on Cu nanolayers (as indicated by arrow), and (d) a deformation twin in an Ag nanolayer that crossed the interface and propagated in the neighboring Cu nanolayers [94].

Fig. 10. (a) Comparison of hardness vs. h-1/2 plots for various fcc/bcc metallic multilayer systems, including Cu/Cr, Cu/Nb, Cu/V and Al/Nb, where h is individual layer thickness [29] and (b) hardness as a function of volume fraction of amorphous layers in C/A multilayers [37].

Table 2 Comparison of peak hardness of different crystal-structured nanolaminates with different h.

Nanolaminates	Crystal structure	Layer thickness, h (nm)	Peak hardness, H_max (GPa)	References
Cu/Nb	fcc/bcc	1.2	7.02	[28]
Cu/Ag	fcc/fcc	2.5	4.5	[99]
Cu/Ni	fcc/fcc	1.75	6.83	[74]
Ag/Ni	fcc/fcc	4	2.38	[98]
Cu/Cr	fcc/bcc	10	6.8	[6]
Cu/V	fcc/bcc	2.5	5.2	[29]
Al/Nb	fcc/bcc	1	4.8	[29]
Cu/Au	fcc/fcc	25	2.45	[128]
Cu/Co	fcc/hcp	1	6	[79]
Cu/W	fcc/bcc	4	8.9	[73]
Cu/Ru	fcc/hcp	5	8.1	[129]
Cu/Zr	fcc/hcp	5	5.8	[130]
Cu/Ta	fcc/bcc	30	7	[131]
Mg/Nb	hcp/bcc	5	2.7	[107]
Mg/Ti	hcp/hcp	2.5	4.2	[38]
Zr/Nb	hcp/bcc	30	5.31	[132]

Fig. 11. (a) Dependence of indentation modulus on h for Cu/Nb nanolaminates (modulus of Cu and Nb films also shown by dashed line for Cu and dotted line for Nb) [33]; (b) cross-section HRTEM image of Cu/W nanolaminates with 8 nm periodicity [73].

Fig. 12. Variation in elastic modulus of: (a) Ag/Cu nanolaminate [71], and (b) Cu/Cr [30] nanolaminate with reduction of modulation period.

Fig. 13. (a) Dependence of ?C on hCu for Cu/Nb and Cu/Zr multilayers with ? = 1.0 and on h for single-layer Nb and Zr films [12], and (b) true stress-strain curves from uniaxial compression tests of Cu/Zr micropillars with five different h, where the average diameter of micropillars is 600 nm [106].

Fig. 14. SEM images of Al/SiC multilayer micropillar deformed due to plastic strain of 8% at: (a) 23 °C and (b) 100 °C [142].

Table 3 Elastic modulus E and hardness H measured from instrumented nano-indentation test in monolithic Al and SiC layers and Al/SiC multilayers at 23 °C and 100 °C temperature [150].

Materials	T =23 °C		T = 100 °C
	E (GPa)	H (GPa)	E (GPa)	H (GPa)
Al	88 ± 6	0.91 ± 0.04	70 ± 3	0.60 ± 0.02
SiC	297 ± 7	35.6 ± 0.7	320 ± 3	30.8 ± 0.8
Al/SiC	128 ± 7	4.9 ± 0.5	118 ± 11	3.7 ± 0.4

Table 4 Effects of h and temperature on the hardness of different Al/SiC nanolaminates [151].

Sample	Hardness (GPa)
	28°C	50°C	100°C	150°C
Al10/SiC50	9.7 ± 0.1	9.2 ± 0.3	7.4 ± 0.1	6.1 ± 0.2
Al25/SiC50	8.1 ± 0.5	7.2 ± 0.2	5.2 ± 0.1	4.0 ± 0.1
Al50/SiC50	5.7 ± 0.3	5.3 ± 0.1	3.9 ± 0.2	3.1 ± 0.3
Al100/SiC50	5.5 ± 0.4	5.2 ± 0.4	4.2 ± 0.3	3.6 ± 0.2
Al50/SiC2	3.3 ± 0.2	2.8 ± 0.1	1.5 ± 0.1	1.2 ± 0.1
Al50/SiC10	4.1 ± 0.1	3.6 ± 0.2	2.6 ± 0.1	2 ± 0.1
Al50/SiC25	4.8 ± 0.1	4.4 ± 0.2	3.5 ± 0.2	2.7 ± 0.2
Al50/SiC100	7.6 ± 0.3	7.2 ± 0.3	5.4 ± 0.1	4.0 ± 0.2

Fig. 15. Schematic of a crack: (a) lying on an interface, (b) on a plane within a layer parallel to the interface, (c) in the edge orientation, and (d) in the surface orientation.

Fig. 16. (a) Schematic of micro-mechanical fracture model, and (b) model predicting ?max at different normalized cohesive strengths ($\tilde{\sigma }$ c) as a function of hCu [12].

Fig. 17. Schematic diagram of deformation mechanism of multilayer strength operating at different length scales [28].

Fig. 18. Schematic of CLS mechanism in nanolaminate structures: glide of dislocation loops in individual layers results in dislocations of deposited interfaces [167].

Fig. 19. Schematic representation of CLS model in A/B multilayer (slip nucleation shown in one layer): (a) dislocation deposited at interface due to glide of dislocation loops confined to layer A (dotted line represents slip plane and dislocation motion is normal to plane of view); (b) introducing CLS activity on two different slip systems, where adjacent segments (labeled 1 and 2) would produce dislocations with Burgers vectors in the interface plane labeled 3 in.(c); (c) new glide loops appear in layer A, experience repulsion from array of edge dislocations formed at interface [28].

Fig. 20. SEM indent images for Ag/Cu multilayers with: (a) λ = 4 nm, (b) λ = 10 nm, and (c) λ = 20 nm [71].

Fig. 21. Indentation-creep depth as a function of load-holding time for Ag/Cu multilayers with: (a) λ = 10 nm, and (b) λ = 50 nm [71].

Fig. 22. SEM images of residual indentation morphologies for: (a) (CuZr)10Cu18, (b) (CuZr)20Cu18, (c) (CuZr)75Cu18 and (d) (CuZr)200Cu18 nanolaminates [37].

Fig. 23. SEM images of post-loaded micropillars for Mg/Ti multilayer specimen with: (a) h = 100 nm, (b) h = 5 nm, and (c) h = 2.5 nm [38].

Fig. 24. SEM cross-sectional view of shear deformation morphologies under microindented Au-Cu multilayers: (a) λ = 25 nm, and (c) λ = 250 nm under load of 250 mN for 5 s, magnified images of (b) and (d) from squared area in (a) and (c) respectively, show different shear deformation characteristics at different length scales [190].

Fig. 25. SEM images of two multilayers: (a) Ni/Ag multilayer, where light layers are Ag and dark layers are Ni; specimen was annealed at 600 °C for 72 h, where breakdown of Ni layers occurred, and (b) Cu/Nb multilayer, where light layers are Nb and dark layers are Cu, specimen was annealed at 800 °C for 3 h, where pinch-off happened in Nb layers [191].

Fig. 26. Transmission electron micrographs showing cross-sectional views of microstructures of 75 nm Cu/Nb multilayers at: (a) 600 °C, and (b) 800 °C temperatures; alignment of grain boundaries along a vertical line is indicated by white arrows in (a), where grooving and shear of layers at triple points are shown in (b) [192].

Fig. 27. Microstructural evolution after annealing in 15 nm Cu-Nb multilayers at: (a) 600 °C, 1 h; and (b) 700 °C, 1 h [193].

Fig. 28. Polycrystalline Ag/Ni multilayers with numerous groove angles, where the groove angle depends on the orientation of each grain at triple junction [191].

Fig. 29. Reflectance (ρ) according to wavelengths (λ) for two different spectrally selective surfaces of Ti-TiNyOx and Cr-Cr2O3 on copper substrate [145].

Table 5 Data of an optimized sample for absorbance and thermal emittance [147].

Optimized sample	Thickness (nm)	Metal content (NC)	Absorbance (α)	Thermal emittance (?)
Top layer	85	0	0.969	-
Middle layer	86	0.005	0.921	-
Base layer	72	0.5	0.763	-
Total	243		0.969	0.047

Fig. 30. (a) Standard Mo-Si multilayers consisting of alternating Mo and Si layers and terminated with an Si layer, and (b) interface-engineered multilayers consisting of Mo and Si layers separated with thin boron carbide layers, where their capping layer presented in this paper is Si [199].

Fig. 31. Reflectance of interface-engineered Mo-Si multilayers at 12.7 and 13.5 nm wavelengths, where measurements were done at 5° off normal. Both consist of 50-bilayer Mo-Si multilayers with boron carbide interfaces [199].

Fig. 32. Microstructures of ARB Cu-Nb composites for: (a) 16 nm at room temperature, (b) 58 nm at room temperature, (c) 16 nm at 450 °C and (d) 58 nm at 450 °C [204].

Table 6 Decrease in corrosion rate with increasing number of layers for CMMA Zn/Ni coating system [216].

Current density (A/dm²)	No. of layers	Potential vs Ag, AgCl / KCL_salt (V)	I_corr (μA/cm²)	Corrosion rate (× 10^-2 mm y^-1)
1.0-4.0	10	-0.914	6.554	13.33
	20	-0.881	4.580	6.791
	60	-1.003	3.067	4.550
	120	-0.823	2.539	3.767
	300	-0.799	1.2225	1.816
	600	-1.176	10.129	15.02
2.0-5.0	10	-0.911	7.555	11.20
	20	-0.956	2.634	3.902
	60	-0.938	1.397	2.073
	120	-0.928	1.153	1.170
	300	-0.942	0.379	0.502
	600	-0.968	4.046	6.00

Fig. 33. SEM images of CMMA (Zn-Ni) coatings: (a) cross-sectional view showing 10 layers, and (b) surface morphology of 4 layered CMMA (Zn/Ni) after corrosion test [216].

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