a dissertation - stanford universitykw371jd4253/... · jedal, mingyang li, michael wiemer, rafael...

GAAS-BASED 1550 NM GAINNASSB LASERS

A DISSERTATION

SUBMITTED TO THE DEPARTMENT OF ELECTRICAL

ENGINEERING

AND THE COMMITTEE ON GRADUATE STUDIES

OF STANFORD UNIVERSITY

IN PARTIAL FULFILLMENT OF THE REQUIREMENTS

FOR THE DEGREE OF

DOCTOR OF PHILOSOPHY

Tomas Sarmiento

March 2013

http://creativecommons.org/licenses/by-nc/3.0/us/

This dissertation is online at: http://purl.stanford.edu/kw371jd4253

© 2013 by Tomas Sarmiento Suarez. All Rights Reserved.

Re-distributed by Stanford University under license with the author.

This work is licensed under a Creative Commons Attribution-Noncommercial 3.0 United States License.

ii



http://purl.stanford.edu/kw371jd4253

I certify that I have read this dissertation and that, in my opinion, it is fully adequatein scope and quality as a dissertation for the degree of Doctor of Philosophy.

James Harris, Primary Adviser


David Miller


Jelena Vuckovic

Approved for the Stanford University Committee on Graduate Studies.

Patricia J. Gumport, Vice Provost Graduate Education

This signature page was generated electronically upon submission of this dissertation in electronic format. An original signed hard copy of the signature page is on file inUniversity Archives.

iii

Abstract

Low-cost, long-wavelength light sources are indispensable for the widespread deploy-

ment of fiber-to-the-home networks. Vertical cavity surface emitting lasers (VCSELs)

are ideal for these applications due to their high fiber-coupling efficiency, low power

consumption, simple packaging and wafer-scale manufacturability. In particular, VC-

SELs emitting in the C-band (1530-1565 nm) are highly desirable given that the fiber

optical loss is minimal in this wavelength range. High-performance 1550 nm InP-

based VCSELs using various distributed Bragg reflector (DBR) technologies have

been demonstrated, but these approaches generally require complex and extensive

processing, leading to high manufacturing costs. In contrast, GaAs-based VCSELs

can be processed in a simple and robust way by exploiting the superior material prop-

erties of Al(Ga)As/GaAs DBRs and the oxidation of AlAs layers for electrical and

optical confinement.

Dilute nitride GaInNAsSb alloys emitting in the 1200-1600 nm wavelength range

can be grown coherently on GaAs substrates, enabling the realization of long wave-

length GaAs-based lasers. Despite significant challenges in the growth of such highly-

mismatched alloys, 1.55 µm GaInNAsSb lasers with relatively low threshold current

densities have been demonstrated. This dissertation describes recent progress on the

development of GaInNAsSb lasers. Optimization of the growth and annealing condi-

tions enabled a four-fold enhancement of the luminescence efficiency of GaInNAsSb

quantum wells with GaNAs barriers. In addition, incorporation of GaAsP barriers

significantly improved the temperature stability of the lasers. The improved quan-

tum wells enabled the realization of the first electrically-pumped GaInNAsSb VCSELs

emitting in the C-band that operate at and above room temperature.

iv

Acknowledgements

First and foremost, I would like to thank my research advisor, Professor James Harris,

for his guidance and support during my doctoral studies. His broad knowledge and

keen intuition were invaluable at different stages of this project.

I would also like to thank Professor David A. B. Miller and Professor Jelena

Vuckovic for serving on the reading committee of this dissertation. I am particularly

grateful to Prof. Vuckovic for the research collaborations with the Harris group and

for the opportunity to join her group as a postdoctoral scholar after graduation. I

would also like to thank Professor Mark Brongersma and Professor Ted Kamins for

being part of my orals committee.

I had the privilege of working closely with Hopil Bae on the development of 1.55

µm GaInNAsSb lasers. Together we spent countless hours doing MBE maintenance

and growing and processing lasers. I would also like to thank Evan Pickett who worked

on this project for a few years, and continue to help me with many Hall measurements

after he left Stanford. I would also like to thank the previous generation of dilute

nitride growers: Seth Bank, Homan Yuen, and Mark Wistey, who, along with Hopil,

trained me in MBE growth and maintenance, material characterization, and laser

processing. I am also grateful to Lynford Goddard, who guided me through laser

testing.

I gratefully acknowledge the contributions of other individuals to this work. Tom

O’Sullivan processed the first generation of VCSELs to perfection. James Ferguson,

Peter Smowton and Peter Blood of Cardiff Unversity characterized the gain medium

of our lasers with the segmented contact method. Robert Kudrawiec and his col-

leagues at Wroclaw University performed optical measurements on our dilute nitrides

v

materials. Fellow MBE grower Angie Lin helped me with TEM characterization.

I am indebted to Angie as well Robert Chen for their help with MBE system

openings and maintenance. I am also grateful to the other MBE growers: Hai Lin,

Yijie Huo, Donghun Choi, Shuang Li, and Sara Harrison, who have helped me with

MBE maintenance. I would also like to thank other members of the Harris group.

Meredith Lee, Paul Lim, Sonny Vo, Dong Liang, Altamash Janjua, Thorsten Hes-

jedal, Mingyang Li, Michael Wiemer, Rafael Aldaz, Vijit Sabnis, Pascale El-Kallassi,

Ed Fei, Yiwen Rong, Ken Leedle, Seonghyun Paik, Larkhoon Leem, Tom Lee and

many others. Special thanks go to the members of the ”grad prep” group: Angie

Lin, Meredith Lee and Paul Lim (Honorary) for their support, encouragement, and

friendship. Many thanks to Gail Chun-Creech for making sure that the group runs

smoothly from an administrative point of view.

I had the opportunity to collaborate with several groups at Stanford. I would

like to thank former and current members of the Vuckovic, Melosh, and Brongersma

groups, including Bryan Ellis, Gary Shambat, Dirk Englund, Ilya Fushman, Yiyang

Gong, Arka Majumdar, Michal Bajcsy, Armand Rundquist, Konstantinos Lagoudakis,

Tom Babinec, Kelly Rivoire, Sonia Buckley, Jared Schwede, Sam Rosenthal, Dan Ri-

ley and Kevin Huang. I learned a great deal from my interactions with them.

Lastly, I would like to thank my family and friends for their continual support.

Even though my father passed away in 1991, he had enough time to inculcate in me

a passion for learning that ultimately led me to pursue this doctoral degree. Finally,

I would like to thank Maria Fernanda—my wife, partner and best friend—for all her

love and support throughout these years.

vi

Contents

Abstract iv

Acknowledgements v

1 Introduction 1

2 MBE Growth of Dilute Nitrides 10

2.1 Molecular Beam Epitaxy . . . . . . . . . . . . . . . . . . . . . . . . . 10

2.2 Plasma-Assisted Epitaxy . . . . . . . . . . . . . . . . . . . . . . . . . 13

2.3 Epitaxy in the low solid solubility limit . . . . . . . . . . . . . . . . . 16

2.4 Surfactant-Mediated Epitaxy . . . . . . . . . . . . . . . . . . . . . . 17

2.5 Annealing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 19

3 Enhancement of luminescence efficiency 22

3.1 Group-V Fluxes Optimization . . . . . . . . . . . . . . . . . . . . . . 22

3.2 Annealing Conditions Optimization . . . . . . . . . . . . . . . . . . . 26

3.3 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 31

4 1550 nm GaInNAsSb Lasers with GaNAs Barriers 33

4.1 Laser Structure and Fabrication . . . . . . . . . . . . . . . . . . . . . 33

4.2 Laser Performance . . . . . . . . . . . . . . . . . . . . . . . . . . . . 35

4.3 Sources of non-radiative recombination . . . . . . . . . . . . . . . . . 37

4.4 Temperature Sensitivity . . . . . . . . . . . . . . . . . . . . . . . . . 42

4.5 On the improvement of GaNAs barriers . . . . . . . . . . . . . . . . . 46

vii

4.6 Conclusion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 51

5 1500 nm GaInNAsSb Lasers with GaAsP Barriers 52

5.1 GaInNAsSb quantum wells with GaAs-GaAsP barriers . . . . . . . . 53

5.2 Laser Performance . . . . . . . . . . . . . . . . . . . . . . . . . . . . 55

5.3 Temperature Sensitivity . . . . . . . . . . . . . . . . . . . . . . . . . 58

5.4 Asymmetric GaNAs/GaAs-GaAsP barriers . . . . . . . . . . . . . . . 59

5.5 Conclusion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 60

6 1550 nm GaInNAsSb VCSELs 62

6.1 Challenges for 1.55 µm AlAs/GaAs DBRs . . . . . . . . . . . . . . . 62

6.2 VCSELs with n- and p-doped DBRs . . . . . . . . . . . . . . . . . . 65

6.2.1 Design . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 65

6.2.2 MBE growth and fabrication . . . . . . . . . . . . . . . . . . . 71

6.2.3 Performance . . . . . . . . . . . . . . . . . . . . . . . . . . . . 75

6.3 Tunnel junction VCSELs . . . . . . . . . . . . . . . . . . . . . . . . . 80

6.3.1 GaAs Tunnel Junctions . . . . . . . . . . . . . . . . . . . . . . 80

6.3.2 Design . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 84

7 Conclusions 88

Bibliography 90

viii

List of Tables

4.1 Monomolecular, radiative and Auger recombination current densities

in the well and the barrier . . . . . . . . . . . . . . . . . . . . . . . . 42

4.2 Characteristic temperature of the threshold current density for ideal

and non-ideal quantum well lasers assuming a dominant recombination

mechanism. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 43

ix

List of Figures

1.1 Estimated Internet data traffic in the United States from 1999 to 2009 2

1.2 Bandwidth requirements to stream various TV formats and capabilites

of access networks . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3

1.3 Maximum transmission distance of various optical communication sys-

tems as a function of bit rate . . . . . . . . . . . . . . . . . . . . . . 4

1.4 Energy gap versus lattice constant for various III-V semiconductors,

including dilute nitride InGaNAs alloys . . . . . . . . . . . . . . . . . 5

1.5 Illustration of the band anticrossing model . . . . . . . . . . . . . . . 7

1.6 Threshold current densities of early InGaNAs quantum well lasers as

a function of lasing wavelength . . . . . . . . . . . . . . . . . . . . . 8

2.1 Schematic of a molecular beam epitaxy chamber . . . . . . . . . . . . 11

2.2 Langmuir probe measurements of a nitrogen plasma . . . . . . . . . . 14

2.3 Nitrogen incorporation as a function of group-III growth rate . . . . . 15

2.4 Saturated ion count as a function of deflection plate voltage . . . . . 16

2.5 PL spectra before and after purification of the nitrogen supplied to the

plasma source . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17

2.6 PL spectra of 1.3-1.6 µm GaInNAs(Sb) quantum wells of various el-

emental compositions, illustrating the benefits of adding antimony to

InGaNAs to extend the emission wavelength . . . . . . . . . . . . . . 18

2.7 Peak PL intensity of GaInNAs(Sb) quantum wells as a function of

antimony flux . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 20

2.8 PL spectra of as-grown and annealed GaInNAsSb quantum wells . . . 21

x

3.1 Peak PL intensity of GaInNAsSb/GaNAs quantum wells grown under

various As/III flux ratios with a fixed Sb flux . . . . . . . . . . . . . 23

3.2 Peak PL intensity of GaInNAsSb/GaNAs quantum wells grown under

various As/III flux ratios with a fixed As/Sb flux ratio . . . . . . . . 24

3.3 Structural evaluation of GaInNAsSb/GaNAs quantum wells grown un-

der various As/III flux ratios and substrate temperatures . . . . . . . 26

3.4 Annealing behavior of GaInNAsSb/GaNAs quantum wells annealed for

60 secs at various temperatures . . . . . . . . . . . . . . . . . . . . . 27

3.5 Peak PL intensity of GaInNAsSb/GaNAs quantum wells annealed un-

der various conditions . . . . . . . . . . . . . . . . . . . . . . . . . . . 28

3.6 Peak PL intensity and emission wavelength of GaInNAsSb/GaNAs

quantum wells annealed in situ . . . . . . . . . . . . . . . . . . . . . 29

3.7 Optimization of the annealing conditions for edge-emitting lasers . . . 30

3.8 PL spectra of GaInNAsSb/GaNAs quantum wells grown and annealed

under various conditions . . . . . . . . . . . . . . . . . . . . . . . . . 31

4.1 Schematic of the edge-emitting laser structure . . . . . . . . . . . . . 34

4.2 Room-temperature, continuous-wave operation of a GaInNAsSb/GaNAs

single quantum well laser . . . . . . . . . . . . . . . . . . . . . . . . . 36

4.3 Pulsed L-I characterisitics of a GaInNAsSb/GaNAs single quantum

well laser . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 36

4.4 Pulsed operation of a GaInNAsSb/GaNAs triple quantum well laser . 37

4.5 Band diagram of the GaInNAsSb/GaNAs quantum well active region 39

4.6 Segmented contact measurements . . . . . . . . . . . . . . . . . . . . 40

4.7 Pulsed L-I characteristics of single and triple GaInNAsSb/GaNAs lasers

as a function of temperature . . . . . . . . . . . . . . . . . . . . . . . 44

4.8 Temperature sensitivity of single and triple GaInNAsSb/GaNAs quan-

tum well lasers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 45

4.9 PL spectra of GaInNAsSb quantum wells with GaNAs and GaAs bar-

riers, illustrating the benefits of using GaNAs barriers . . . . . . . . . 47

xi

4.10 Peak PL intensity of GaInNAsSb/GaNAs quantum wells a a function

of barrier thickness . . . . . . . . . . . . . . . . . . . . . . . . . . . . 48

4.11 Pulsed L-I characteristics of GaInNAsSb/GaNAs lasers with various

barrier thicknesses . . . . . . . . . . . . . . . . . . . . . . . . . . . . 49

4.12 PL spectra of GaInNAsSb quantum wells surrounded by GaNAs bar-

riers grown with various As/III flux ratios. . . . . . . . . . . . . . . . 50

5.1 Energy gap versus lattice constant for various III-V semiconductors . 53

5.2 Band diagram of a GaInNAsSb quantum well surrounded by GaAs-

GaAsP barriers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 54

5.3 Structural evalution of GaInNAsSb with GaAs-GaAsP barriers . . . . 55

5.4 PL spectra of GaInNAsSb quantum wells with GaAs-GaAsP and GaNAs

barriers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 56

5.5 Pulsed L-I characteristics of lasers with GaAs-GaAsP and GaNAs bar-

riers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 57

5.6 TEM micrograph of a GaInNAsSb quantum well surrounded by GaAs-

GaAsP barriers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 57

5.7 SIMS depth profile of a GaInNAsSb quantum well with GaAs-GaAsP

barriers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 58

5.8 Temperature sensitivity of GaInNAsSb lasers with GaNAs and GaAs-

GaAsP barriers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 59

5.9 Structural evaluation of GaInNAsSb quantum wells with asymmetric

barriers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 60

5.10 PL spectrs of GaInNAsSb quantum wells with asymmetric barriers . 61

6.1 Absoprtion coefficient of p-doped GaAs and refractive index contrast

of the AlAs-GaAs material system as a function of wavelength . . . . 63

6.2 Reflectivity of undoped and p-doped AlAs/GaAs DBRs as a function

of mirror pairs at various wavelengths . . . . . . . . . . . . . . . . . . 64

6.3 Calculated VCSEL threshold current density as a function of top DBR

reflectivity . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 67

xii

6.4 Grading and doping profiles of the designed p-doped AlGaAs/GaAs

DBR . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 69

6.5 Scattering loss as a function of aperture diameter . . . . . . . . . . . 71

6.6 Schematic of the 1.53 µm GaInNAsSb VCSEL structure . . . . . . . 72

6.7 XRD scan of the active region of the VCSELs . . . . . . . . . . . . . 73

6.8 Optimization of the ex-situ annealing conditions for various in-situ

annealings . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 74

6.9 Measured reflectivity spectrum of the VCSEL structure . . . . . . . . 75

6.10 Pulsed operation of a 1.53µm GaInNAsSb VCSEL . . . . . . . . . . . 76

6.11 Continuous-wave operation of a 1.53 µm GaInNAsSb VCSEL . . . . . 76

6.12 Measurements to determine the thermal resistance of the VCSEL . . 77

6.13 Thermal resistance of the VCSEL as a function of aperture diameter . 79

6.14 Schematic of an intracavity-contacted VCSEL structre . . . . . . . . 80

6.15 Electron concentration in a Si-doped GaAs layer as a function of silicon

cell temperature . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 81

6.16 Growth temperature dependence of the electron concentration in a Si-

doped GaAs layer . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 82

6.17 Band structure of a GaAs tunnel junction . . . . . . . . . . . . . . . 83

6.18 Measured current-voltage characteristics of a GaAs tunnel junction

with ∼1.2×1019 cm−3 active donors in the n++ layer. . . . . . . . . . 84

6.19 Measured current-voltage characteristics of a GaAs tunnel junction

with ∼4.5×1018 cm−3 active donors in the n++ layer. . . . . . . . . . 85

6.20 Refractive index profile along with the electric field intensity in the

designed tunnel-junction VCSEL structure . . . . . . . . . . . . . . . 85

6.21 Schematic of the designed 1.55 µm GaInNAsSb tunnel-junction VCSEL 86

6.22 Reflectivity spectrum of a GaInNAsSb tunnel-junction VCSEL . . . . 87

xiii

Chapter 1

Introduction

Internet data traffic is ever increasing. As an example, Fig. 1.1 shows the Internet data

traffic in the United States from 1999 to 2009 as estimated by the Minessota Internet

Traffic Studies (MINTS) [1]. Even after the burst of the dot-com and telecommuni-

cation bubbles, data traffic has increased steadily with an annual growth rate of 54%.

This growth rate is consistent with other studies that indicate that the traffic in the

United States roughly doubles every two years [2]. It is estimated that over 50% of

this traffic is related to real-time entertainment services.

Video on demand of high definition TV (HDTV) is one of these services that is

driving the demand for bandwidth in the access networks. Streaming HDTV requires

download speeds between 5 and 15 Mbps, and these bandwidth requirements are

going to be pushed to unprecedented levels when the next generation of ultra HDTV

formats are introduced. The 4K TV format with four times the spatial resolution

of HDTV requires download speeds around 50 Mbps while the 8K TV format with

sixteen times the spatial resolution of HDTV requires around 200 Mbps.

The access networks that connect the end user to the optical backbone of the

Internet network are typically cable or digital subscriber line (DSL) networks. Cable

networks that use the DOCSIS2.0 standard provide download speeds up to 38 Mbps.

DSL networks support download speeds up to 24 Mbps with the asymmetric DSL

(ADSL) format and up to 55 Mbps with the very-high-bit-rate DSL (VDSL) format.

The capabilities of the access networks and the bandwidth requirements to stream

1

CHAPTER 1. INTRODUCTION 2

2000 2002 2004 2006 2008 2010101

102

103

104

Year

Dat

a Tr

affic

(P

etaB

ytes

/mon

th)

54% annual growth

MINTS data (mid-range)

Figure 1.1: Mid-range of the Internet data traffic in the United States from 1999 to2009 as estimated by the Minessota Internet Traffic Studies (MINTS).

the different TV formats are summarized in Fig. 1.2. Clearly, the download speeds

supported by these access networks are sufficient to stream HDTV but inadequate to

stream ultra HDTV.

In contrast to cable and DSL networks, fiber-to-the-home (FTTH) networks can

support sustained bandwidths in excess of 1 Gbps. These high bandwiths are achieved

by extending the optical network all the way to the end user. The high cost of these

networks has limited their deployment in North America. In some cases, the cost of

the optical components—in particular of the light source— represents a large fraction

of the total cost of the network. Thus, reducing the cost of the light source would

enable widespread deployment of high-bandwidth FTTH networks.

Ideally, these light sources should emit in the C-band, from 1530 to 1565 nm, as

the optical fiber loss is minimal in this wavelength range. In addition, the material

dispersion in the optical fiber can be engineered to be minimal in this range. Minimum

attenuation and minimum dispersion ensures maximum transmission distances at high

speeds as can be seen in Fig. 1.3.

Two types of ligth sources can be used in optical networks: distributed-feedback

(DFB) lasers and vertical cavity surface emitting lasers (VCSELs). DFB lasers emit


0 50 100 150 200Download data rate (Mbps)

Cap

abili

ties

Req

uire

men

ts SDTV

HDTV

4K TV

8K TV

ADSL2+

Cable (DOCSIS2.0)

VDSL

Figure 1.2: Bandwidth requirements to stream various TV formats and capabilites ofcable and DSL access networks.

a highly asymmetric elliptical beam, and hence a cylindrical lens and precise align-

ment are required to couple this beam into an optical fiber, which complicates device

packaging and increases the cost of the laser. The cost of DFB lasers is acceptable

for long-haul optical networks but their cost is too high for FTTH networks. On

the other hand, VCSELs have several properties that enable their manufacture at

lower cost. The emitted beam has a cylindrical profile and hence the optical power

can be coupled efficiently into optical fibers, which allows simple, lens-free packaging.

Furthermore, VCSELs can be manufactured and tested at the wafer scale such that

only the devices that meet the required specifications are diced and packaged, which

reduces the cost significantly. VCSELs emitting at 1.55 µm are therefore the ideal

low cost light sources for FTTH networks.

In a VCSEL, the round-trip gain is small and hence, distributed Bragg reflec-

tors (DBRs) with very high reflectivities are required. This is a major difficulty for

InP-based VCSELs as the refractive index contrast of the alloys that can be grown

coherently on InP substrates is low, with the notable exception of AsSb-based DBRs.

Therefore, a large number of mirror pairs are required to achieve suitable reflectivities.


0.001 0.01 0.1 1 10 100

10

100

1000

Bit Rate (Gbps)

Dis

tanc

e (k

m)

850 nm

1.3 µm

1.55 µm Dispersion shifted fiber

Figure 1.3: Maximum transmission distance of various optical communication systemsas a function of bit rate.

Moreover, InP-based DBRs invariably incorporate a quaternary alloy. Due to alloy-

disorder scattering of phonons, quaternary alloys have very low thermal conductivi-

ties [3], which hinders heat removal from the active region, a serious complication as

the GaInAsP/InP active regions are temperature sensitive. Despite these difficulties,

1.55 µm VCSELs have been demonstrated using different DBR technologies, includ-

ing all epitaxial AlGaAsSb/AlAsSb [4] and InAlGaAs/InAlAs [5] DBRs, InP/air-gap

DBRs [6], wafer fused DBRs [7, 8], and a combination of epitaxial DBRs with di-

electric DBRs [9, 10], metamorphic DBRs [11, 12] or high contrast grating (HCG)

mirrors [13]. While these VCSELs exhibit high performance, they require complex

and extensive processing that leads to high manufacturing costs.

In contrast, GaAs-based VCSELs can be manufactured at extremely low cost

by exploiting the superior properties of GaAs-based materials. AlAs is nearly lat-

tice matched to GaAs and can be used to form AlAs/GaAs DBRs with high index

contrast. The higher index contrast reduces the number of mirror pairs required

to achieve suitable reflectivities. And, compared to InP-based DBRs, AlAs/GaAs

DBRs have higher thermal conductivities since these DBRs are composed of binary

alloys, rather than ternary and quaternary alloys. The higher thermal conductivities


5.55 5.65 5.75 5.85 5.95 6.05

0.4

0.6

0.8

1.0

1.2

1.4

Lattice Constant (A)

Ene

rgy

Gap

(eV

)

GaAsGaNyAs1-y

InAs

InxGa1-xAs

InP

InxGa1-xNyAs1-y

1.31 µm

1.55 µm

Figure 1.4: Bandgap energy versus lattice constant for various III-V semiconduc-tors, including dilute nitride InxGa1−xNyAs1−y (y ≤ 4%) alloys that can be growncoherently on GaAs substrates.

and thinner DBR stacks result in more efficient heat removal capabilities, improving

the performance and maximum temperature of operation of the VCSELs. Moreover,

AlAs can be selectively oxidized to define a current aperture, which simplifies VCSEL

processing as it eliminates the regrowth, ion implantation, or critical etching steps

necessary to define the current aperture in InP-based VCSELs.

While AlAs/GaAs DBRs are ideal for VCSELs, the realization of a GaAs-based

active region that emits at 1.55 µm is extremely challenging. These difficulties are

illustrated in Fig. 1.4. Conventional GaAs-based VCSELs that emit at 980 nm use

strained InGaAs (In∼20%) quantum wells. Increasing the indium concentration of

the quantum wells extends the emission wavelength but increases the lattice constant.

Thus, strain limits the indium concentration and the emission wavelength to ∼1100

nm. However, if dilute amounts of nitrogen are added to InGaAs, to form InGaNAs,

both the lattice constant and the band gap of the alloy decreases. Hence, InGaNAs

quantum wells emitting at 1.3 and 1.55 µm can be grown coherently on GaAs, en-

abling the realization of long-wavenlength GaAs-based lasers. This was recognized

by Kondow et al. [14,15] based on the discovery of the anomalous bandgap reduction


of GaAs with the addition of nitrogen [16].

The anomalous reduction in the energy gap with the addition of dilute amounts

of nitrogen can be explained with the semiempirical band anti-crossing (BAC) model

[17]. Due to the large electronegativity of nitrogen, the substitution of a few percent of

the host anions with nitrogen results in a considerable perturbation to the electronic

potential surrounding the nitrogen atoms. Since this perturbation is localized in real

space it forms an energy level extended in momentum space, which is located above

the conduction band minimum of the host semiconductor. The repulsive interaction

of the nitrogen energy level and the conduction band results in the formation of two

energy subbands, E+ and E−, with energy dispersion relations described by

E±(k) =EN + EM(k)±

√(EN − EM(k))2 + 4C2

NMx

2(1.1)

where EN is the energy of the nitrogen level1, EM(k) is the parabolic dispersion rela-

tion for the conduction band of the host semiconductor, CMN is a coupling constant

describing the strength of the interaction and x is the nitrogen concentration. The

conduction subbands are located above (E+) and below (E−) the conduction band

of the host semiconductor as can be seen in Figure 1.5. The formation of the lower

conduction subband E− causes the reduction of the bandgap energy. Compared to the

host conduction band, this conduction subband is flatter near its minimum, leading

to larger electron effective mass.

Besides being nearly lattice matched to GaAs, InGaNAs lasers have additional

advantages over their InGaAsP counterparts. Compared to InGaAsP/InP quantum

wells, InGaNAs/GaAs quantum wells have larger conduction band offsets [15]. The

larger offset and the heavier electron effective mass results in better electron confine-

ment and hence lower temperature sensitivity.

InGaNAs lasers emitting in the 1.3 µm range with low threshold current densities

[18–22] and high performance VCSELs [23–25] have been demonstrated. But as more

nitrogen is added to the InGaNAs quantum wells to extend the emission wavelength

to the 1.55 µm range, the optical quality of the material degrades substantially. This

1Relative to the top of the valence band of the host semiconductor.


BandAnticrossing(repulsion)

E+

E-

ENitrogen

EC

hh

lhso

k

E

BandAnticrossing(repulsion)

E+

E-

ENitrogen

EC

hh

lhso

k

E

Figure 1.5: Illustration of the band anticrossing model, showing the splitting of theconduction band into two non-parabolic subbands E+(k) and E−(k).

is illustrated in Fig. 1.6 that shows the increase in laser thresholds as more nitrogen is

added to the quantum wells to extend the lasing wavelength [26]. This increase is due

to the low solid solubility of nitrogen in InGaAs, which leads to phase segregation and

roughening that degrades the optical quality of the material. The propensity of the

alloy to phase segregate increases as more indium and nitrogen are added to the alloy.

To improve the laser performance, it was necessary to suppress the phase segregation.

This can be accomplished by growing at very low substrate temperatures [27].

Alternatively, phase segregation can be suppresed by adding antimony to form

GaInNAsSb alloys. The optical quality of GaInNAsSb was further improved through

the optimization of the growth and plasma conditions, and the use of strain compen-

sating GaNAs barriers. These improvements led to the demonstration of lasers with

threshold current densities in the 300 A/cm2 range [28].

While these laser thresholds are relavitely low, there are still complex challenges

facing the development of 1.55 µm GaInNAsSb lasers. The optical quality of the

material needs to be improved as there is a still a significant concentration of defects.


1.20 1.25 1.30 1.35 1.40 1.45 1.500.0

0.5

1.0

1.5

2.0

2.5

3.0

3.5

0.0

0.5

1.0

1.5

2.0

2.5

Thre

shol

d Cu

rrent

Den

sity

(kA/

cm2 )

Wavelength (µm)

Nitr

ogen

Con

tent

(%)

InP Lasers*

1.20 1.25 1.30 1.35 1.40 1.45 1.500.0

0.5

1.0

1.5

2.0

2.5

3.0

3.5

Thre

shol

d Cu

rrent

Den

sity

(kA/

cm2 )

Wavelength (µm)

“NitrogenPenalty”

Figure 1.6: Threshold current densities of early InGaNAs quantum well lasers as afunction of lasing wavelength.

The temperature stability of these lasers is poor, which means that the threshold

current density increases considerably as temperature increases. This is caused by

the small conduction and valence band offsets between the GaNAs barriers and the

quantum well, which allows the carrier to escape thermionically from the quantum

well into the barriers. There the carriers recombine non-radiatively due to the poor

optical quality of GaNAs.

Regarding VCSEL develoment, electrically-pumped GaInNAsSb VCSELs emit-

ting beyond 1.31 µm have been demonstrated but their operation was limited to

sub-zero temperatures [29,30].

This work had three main objectives: the enhancement of the optical quality of

GaInNAsSb quantum wells, the improvement of the temperature sensitivity of 1.55

µm GaInNAsSb edge-emitting lasers and the realization of GaAs-based 1.55 µm GaIn-

NAsSb VCSELs that operate at room temperature. This dissertation is organized as

follows. Chapter 2 presents a brief introduction to the molecular beam epitaxy (MBE)

growth and describes techniques necessary for the growth of dilute nitrides, such as

plasma-assisted and surfactant-mediated epitaxy. Chapter 3 describes improvements


in the optical quality of GaInNAsSb/GaNAs quantum wells through the optimiza-

tion of the growth and annealing conditions. The performance of the lasers based on

these improved quantum wells is presented in Chapter 4. The sources of non-radiative

recombination in these lasers are identified and strategies to improve the laser per-

formance are discussed. Efforts to improve the temperature stability of GaInNAsSb

lasers using strain-compensating GaAsP barriers are presented in Chapter 5. The

realization of the first electrically-injected GaAs-based VCSELs emitting beyond 1.4

µm that operate at room temperature is presented in Chapter 6. In addition, a new

design to improve the performance of the VCSELs is proposed. Chapter 7 presents a

summary of this work and suggestions for future work.

Chapter 2

MBE Growth of Dilute Nitrides

Molecular beam epitaxy was used to grow the quantum well and laser structures in

this work. This chapter provides a brief introduction to MBE as well as the basics of

plasma assisted and surfactant mediated epitaxy, which are necessary to address the

challenges imposed by the growth of dilute nitride materials.

2.1 Molecular Beam Epitaxy

In molecular beam epitaxy, the growth of epitaxial films occurs in a ultra-high vac-

uum (UHV) enviroment through the reaction of molecular or atomic beams with a

heated substrate. These beams are produced by evaporation or sublimation of high-

purity elemental sources and are transported towards the substrate with negligible

interactions due to the long mean free paths provided by the UHV environment. The

substrate is heated to provide thermal energy for the incident atoms to find lattice

sites. Since the growth is performed under UHV (Base pressure ∼10−10 Torr), there

is minimal incorporation of impurities into the growing film.

A schematic of a MBE system is shown in Fig. 2.1. Effusion or Knudsen cells

contain the elemental sources in inert crucibles that are heated raditively. The tem-

perature of the cell controls the rate of evaporation or sublimation. Mechanical shut-

ters provide on/off modulation of the molecular beams. The use of shutters combined

with the slow growth rate provide atomic-layer control of thickness and composition.

10

CHAPTER 2. MBE GROWTH OF DILUTE NITRIDES 11

Effusion Cells

Effusion Cells

ShuttersShutters

RHEED Gun

RHEED Gun

Substrate Holder & Heater

Substrate Holder & Heater

Fluorescent Screen

Fluorescent Screen

View PortsView Ports

Cryoshroud

Gate Valve to Buffer

Chamber

Gate Valve to Buffer

Chamber

Ion GaugeIon Gauge

SampleSample

Figure 2.1: Schematic of a molecular beam epitaxy chamber.

The growth processes can be monitored in situ using reflection high-energy electron

diffraction (RHEED).

One important aspect of MBE is that the growth rate is independent of the sub-

strate temperature. This is in contrast to metal organic chemical vapor deposition

(MOCVD), where the dissociation efficiency of the precursors is strongly dependent

on the substrate temperature and hence the growth rate depends on the substrate

temperature. Therefore the substrate temperature for MBE is not limited by the

minimum temperature at which the precurssor dissociates and low temperatures can

be used. The kinetics of the growth processes are controllled solely by the substrate

temperature.

Stiochiometric growth of III-V semiconductors is straightforward in MBE. At the

growth temperatures used, all the group-III atoms incident on the substrate stick and

are incorporated, and only those group-V atoms required to bond to group-III atoms

incorporate [31]. Therefore, the growth rate is solely controlled by the group-III fluxes

while stiochiometry is achieved simply by growing in an excess flux (or overpressure)

of group-V elements.


System Description

Two MBE chambers connected under UHV were used for this work. One system is

devoted to the growth of dilute nitride alloys while the other is used for the growth of

AlAs/GaAs DBRs for VCSELs and AlGaAs cladding layers for edge-emitting lasers.

Gallium and indium were supplied by dual filament 400 g SUMO [32] effussion cells.

These cells were operated in ’hot-lip’ mode to prevent condensation of material at the

orifice of the crucible, and thereby reduce oval defects [33]. Aluminum was supplied

by either a dual-filament cold lip 400 g SUMO source or a single-filament 60 cc cell

with a conical crucible. The aluminum SUMO cell was operated in ’cold-lip’ mode

to prevent aluminum from creeping beyond the crucible orifice. The tip filament was

left unheated or set to 80 ◦C below the base temperature for improved flux uniformity

across the wafer. A 500 cc arsenic valved cracker with the cracker zone temperature

set to 850 ◦C supplied dimeric arsenic. Antimony was supplied by a 175 cc unvalved

cracker. The cracker temperature is set to 850 ◦C, which produces mostly monomeric

antimony at the beam fluxes used (10−7 Torr) [34].

Phosphorous was supplied by a GaP decomposition source, which is a conventional

effusion cell with a Ga-trapping cap that removes gallium atoms from the phosphorous

beam. Phosphorous sublimates from GaP as P2 with minimal generation of P4.

This is important because P4 condenses as white phosphorous, which is flammable

upon exposure to oxygen and therefore complicates system maintenance. The major

drawback of this type of source is that the phosphorous flux cannot be switched

on/off instantly because phosphorous escapes around the closed shutter due to its

high vapor pressure. The phosphorous flux is only reduced by a factor of 5 when

the shutter is closed. Thus, in order to grow structures with abruptly changing

phosphorous concentrations, the growth has to be interrupted and the temperature

of the source has to be changed. This problem can be solved using a phosphorous

valved cracker.

Silicon and beryllium—used as n-type and p-type dopants, respectively—were

supplied by 5 cc effusion cells via sublimation. Carbon is also a p-type dopant and

was supplied by a CBr4 source in the AlAs/GaAs system.

The source used to supply nitrogen is discussed in detail in the next section.


2.2 Plasma-Assisted Epitaxy

Generation of reactive nitrogen

Due to the strength of the nitrogen-nitrogen bond and the high vapor pressure of

nitrogen, a plasma source is required in solid-source MBE to dissociate the nitro-

gen molecule and generate reactive nitrogen. This plasma also generates ions that

cause significant damage to the surface. To minimize ion damage, radiofrequency (rf)

plasma sources are used as these plasmas generate fewer ions and have higher atomic

dissociation yields than do electron cyclotron resonance (ECR) and DC plasmas [35].

The nitrogen source used for this work is an rf plasma source manufactured by

SVT Associates. In this source, the plasma is ignited by an rf field applied through

a water-cooled copper coil, which is wrapped around a pyrolitic boron nitride (pBN)

crucible that confines the nitrogen gas. High purity (99.999%) nitrogen is introduced

into the crucible with a leak valve and reactive escape through holes in the aperture

of the crucible. The crucible is of the ”uni-bulb” design and has 4 holes arranged in

a Y pattern. The nominal hole diameter is 0.2 mm. This crucible is different from

that previously used which consisted of a tube and a separate aperture plate. The

integrated crucible provides better plasma stability and reproducibility.

The plasma is operated in the inductively-coupled (high-brightness) mode for

efficient generation of reactive nitrogen. The plasma properties are dictated by the

applied rf power and the pressure inside the crucible, which is controlled by the N2 flow

rate. These parameters must be optimized to balance the competing requirements

of high dissociation efficiency and low ion counts and ion energies. On one hand,

increasing the rf power increases the dissociation efficiency but increases the number

of ions and their energies. On the other hand, increasing the pressure decreases the

ion counts and ion energies but decreases the dissociation efficiency [36, 37]. The

parameter range is limited by the stability of the plasma. The higher the N2 flow rate

the higher the rf power required to produce a stable plasma.

To obtain a plasma with adqueate dissociation efficiencies and low ion counts and

energies, a moderate rf power (300 W) was applied and the N2 flow rate was set to

the highest value that results in an stable plasma. The benefits of this approach can


-20 0 20 40 60

-2

-1

0

1

2

Probe Voltage (V)

Pro

be C

urre

nt (n

A)

33 eV

28 eV

0.095 sccm0.115 sccm 0 20 40 60

-0.3

-0.2

-0.1

0

Voltage (V)

dI/d

V (a

. u.)

0.095 sccm0.115 sccm

Figure 2.2: Langmuir probe measurements of a nitrogen plasma at two differentnitrogen flow rates. The inset show the derivative of the measurement. For thesemeasurements the grid of the beam flux gauge was used as the remote Langmuirprobe.

be quantified using the beam flux gauge as a remote Langmuir probe [38,39]. In this

experiment, a DC voltage is applied to the filament of the beam flux gauge and the

current collected is measured by a picoammeter. Applying a negative voltage, attracts

ions to the probe and hence the measured probe current is a relative measurement

of the number of ions generated by the plasma. If a positive bias is applied to the

probe, ions with energies less than the applied voltage are repelled and the measured

current is due to electrons and ions with energies greater than the applied voltage. An

estimate of he maximum ion energy can be obtained from the probe current-voltage

characteristics at positive biases. As an example, Figure 2.2 shows the Langmuir

measurements of a nitrogen plasma operated at 300 W with flow rates of 0.095 and

0.115 sccm. A ∼17% reduction in the N2 flow rate resulted in a ∼30% increase in

the number of ions generated by the plasma and increased the maximum ion energy

from 28 to 33 eV.

To operate the plasma at optimal conditions, the nitrogen concentration was varied

by changing the group-III growth rate, instead of changing the N2 flow rate. This is


Figure 2.3: Nitrogen concentration in dilute nitride GaNAs films as a function of thegroup-III growth rate.

straigthforward as the nitrogen incorporation in dilute nitride arsenides is inversely

proportional to the group-III growth rate [40]. This is illustrated in Fig. 2.3 that

shows the nitrogen concentration in dilute nitride GaNAs layers as a function of the

group-III growth rate.

The ions produced by the plasma cause significant damage to the surface. This

damage can be reduced by using parallel metal plates at the aperture of the cell

that are biased such that the electric field across them deflects ions away from the

surface [41]. Typically, ±50 V is sufficient to deflect most of the ions as is evident in

Fig. 2.4 that shows the relative ion flux measured by the Langmuir probe technique

as a function of deflection plate voltage. Unless otherwise noted, all the samples in

this work were grown using deflection plates biased to -40 V.

Purification of nitrogen is critical to obtain dilute nitrides of the highest optical

quality. The nitrogen supplied was 99.999% and filtered with a part-per-billion pu-

rifier. In addition, the nitrogen supply lines were baked after exposure to air and

care was taken to eliminate any leaks in the lines. The importance of these steps is

illustrated in Fig. 2.5 that shows a substantial improvement in luminescence efficiency


-90 -60 -30 0 30 60 900

1

2

3

4

5

6

Deflection Plate Voltage (V)

Sat

urat

ed Io

n C

urre

nt (

nA)

Figure 2.4: Saturated ion current as a function of deflection plate voltage. For thesemeasurements the filament of the beam flux gauge was used as the remote Langmuirprobe.

after purification.

2.3 Epitaxy in the low solid solubility limit

In addition to the complications associated with the generation of reactive nitro-

gen, there are some challenges for the growth of dilute nitride (In)GaNAs alloys.

These challenges stem from the low solid solubility of nitrogen in (In)GaAs [42, 43].

Differences in the crystal structure of the binary (or ternary) constituents create a

miscibililty gap. Hence during growth, InGaNAs tends to phase separate into regions

of InGaAs and InGaN. This phase segregation leads to roughening of the surface and

severe degradation of the material quality.

To prevent phase segregation, dilute nitrides are grown at low substrate temper-

atures to limit the thermal energy disponible for this process to occur. However, at

these low growth temperatures a significant concentration of point defects is incor-

porated into the material, degrading its optical quality. The concentration of defects

increases with decreasing growth temperature. Therefore, there is an optimal growth


Figure 2.5: PL spectra showing improvement in luminescence efficiency after purifi-cation of the nitrogen gas injected into the plasma source.

temperature that balances phase segregation and point defect formation. The optical

quality of the material degrades due to phase segregation if higher growth temper-

atures are used or due to point defect formation if lower growth temperatures are

used.

The propensity of InGaNAs to phase segregate increases as the indium or nitro-

gen concentration increases and therefore lower growth temperatures are required

for the growth of InGaNAs with the relatively high indium and nitrogen concentra-

tions required to achieve emission at 1.55 µm. This narrows the growth temperature

window considerably as the onset of phase segregation occurs at lower growth tem-

peratures. The growth temperature window can be extended by using a surfactant

that suppresses phase segregation, allowing higher growth temperatures. The use of

surfactants for the growth of dilute nitrides is discussed in the next section.

2.4 Surfactant-Mediated Epitaxy

A surfactant is a surface segregating species that modifies the epitaxial growth kinet-

ics. Depending on its effect on the surface diffusion length, the surfactant is classified


0.75 0.80 0.85 0.90 0.95 1.00

100

101

102

103

1600 1500 1400 1300

More In, Sb, N

GaInNAsSb

GaInNAs

GaInNAs

PL In

tens

ity (a

.u.)

Energy(eV)

Wavelength (nm)

More In and N

Figure 2.6: PL spectra of 1.3-1.6 µm GaInNAs(Sb) quantum wells of various elementalcompositions, illustrating the benefits of adding antimony to InGaNAs to extend theemission wavelength.

as non-reactive or reactive [44, 45]. A non-reactive surfactant bonds weakly with the

constituent atoms of the crystal such that an adatom can use the additional bonds

provided by the surfactant to facilitate its migration on the surface as less energy

is required to break the surfactant-adatom bond. Hence, non-reactive surfactants

enhance the surface diffusion length. In contrast, reactive surfactants decrease the

surface diffusion length. The bond between the surfactant and the constituent atoms

is stronger and hence more energy is required to break that bond. The surfactant and

the adatom can exchange places such that the adatom is incorporated underneath the

surfactant layer. The effects of reactive surfactants are beneficial for the growth of

strained heterostructures that tend to form islands as a mechanism to relieve strain.

Antimony acts as a reactive surfactant during the growth of InGaNAs [46–48], re-

ducing the surface diffusion length of the impinging adatoms, and consequently, sup-

pressing phase segregation and roughening. This allows higher growth temperatures

and therefore improved material quality as fewer point defects are incorporated at

higher growth temperatures. Since antimony is a reactive surfactant, it incorporates

into the alloy, redshifting its emission wavelength. The benefits of using antimony are

illustrated in Fig. 2.6 showing the PL spectra of 1.3-1.6 µm GaInNAs(Sb) quantum


wells with various elemental compositions. If more indium and nitrogen are added

to InGaNAs quantum wells to extend the emission wavelength beyond 1.3 µm, the

material quality degrades due to phase segregation. If antimony is added instead, the

optical quality of the quantum wells improves and its emission wavelength redshifts.

Antimony allows the addition of more indium and nitrogen to extend the emission

wavelength to the 1.55 µm range.

The presence of antimony on the surface influences the incorporation of other

elements in the alloy. Nitrogen incorporation in enhanced significantly, up to 50%

[49,50]. The indium concentration can be reduced as both antimony and indium are

large atoms that compete for incorporation [51].

The antimony flux used during the growth of GaInNAsSb quantum wells must

be selected to obtain an optimal surface diffusion length [51, 52]. Low antimony

fluxes lead to high surface diffusion lengths that are ineffective at suppresing phase

segregation. High antimony fluxes result in low surface diffusion lengths that cause

incorporation of defects. These effects are illustrated in Fig. 2.7 that shows the

peak PL intensity of GaInNAsSb quantum wells grown at various antimony fluxes.

There is an optimal antimony flux that yields that highest luminescence efficiencies.

Growing at lower or higher fluxes degrades the optimal quality of the material. The

optimal antimony flux depends on the indium and nitrogen concentrations and on

the substrate temperature and growth rate. Higher growth temperatures and lower

growth rates lead to longer surface diffusion lengths and hence the optimal antimony

flux is expected to be increase with increasing growth temperatures and decreasing

growth rates.

2.5 Annealing

The optical quality of as-grown dilute nitrides is extremely poor due to the large

concentration of defects that incorporate during growth. Thermal annealing dramat-

ically improve the optical quality by removing some of these defects [53–55].Figure 2.8

shows the PL spectra of as-grown and annealed 1.55 µm GaInNAsSb quantum wells.

The luminescence efficiency of the quantum wells is improved by a factor of 20 upon


0 2 4 6 8 100

0.2

0.4

0.6

0.8

1.0

Antimony Flux (10 -8 Torr)

Pea

k P

L In

tens

ity (

a. u

.)

Figure 2.7: Peak PL intensity of GaInNAs(Sb) quantum wells as a function of anti-mony flux [28].

annealing.

Interestingly, the emission wavelength of dilute nitrides blueshifts upon anneal-

ing. For GaInNAs(Sb) alloys this blueshift is due to a rearrangement of the nitrogen

nearest neighbors [56–58]. The most favorable nearest neighbor configuration is the

one that minimizes the free energy of the alloy. The cohesive bond energy and lo-

cal strain energy contribute to this free energy. But during growth the local strain

does not play an important role as the strain can be relieved towards the surface.

Since the Ga-N bond is stronger than the In-N bond1, during growth nitrogen atoms

bond preferentially to gallium atoms to minimize the cohesive bond energy. During

annealing, however, the local strain energy also contributes to the free energy. The

local strain energy is lower for In-N bonds as the In-N bond is longer and hence less

stretched than the Ga-N bond. Even though the cohesive bond energy increases,

the local strain and total energies decrease as the nitrogen atoms replace their gal-

lium neighbors with indium atoms [60]. Therefore, during annealing, nitrogen tends

to move from a gallium-rich to an indium-rich environment. The different nitrogen

1The Ga-N and In-N cohesive bond energies are 2.24 and 1.93 eV/bond, respectively [59].


1450 1500 1550 1600 1650 17000

0.2

0.4

0.6

0.8

1.0

Wavelength (nm)

PL

Inte

nsity

(a.

u.)

10X

AnnealedAs-grown

Figure 2.8: PL spectra of as-grown and annealed GaInNAsSb quantum wells.

nearest neighbor environment results in different bandgaps, leading to the observed

blueshift [61]. The magnitude of the blueshift increases with indium concentration as

the number of indium-rich sites increases.

The annealing procedure is thus critical as it strongly affects the optical quality

and the emission wavelength of the dilute nitrides.

Chapter 3

Enhancement of luminescence

efficiency

This chapter presents the optimization of growth and annealing conditions that led

to a substantial improvement of the luminescence efficiency of GaInNAsSb quantum

wells with GaNAs barriers emitting in the 1.55 µm range.

3.1 Group-V Fluxes Optimization

In this section, the dependence of the optical quality of GaInNAsSb/GaNAs quantum

wells on the arsenic and antimony fluxes is investigated.

Based on thermodynamic considerations, it can be concluded that low substrate

temperatures and high arsenic and nitrogen fluxes are required to prevent phase

segregation during the epitaxial growth of dilute nitride GaNAs alloys [62]. But, at

these low growth temperatures, the use of high arsenic fluxes leads to the introduction

of arsenic-associated point defects, such as arsenic antisites and gallium vacancies,

that degrade the optical quality of (In)GaNAs. Therefore, the optimal As flux for

the growth of (In)GaNAs must be high enough to avoid phase segregation, but low

enough to minimize the incorporation of arsenic-related point defects.

The tendency of InGaNAs alloys to phase segregate increases as the indium and/or

nitrogen concentration increase. Therefore, very low substrate temperatures (≤390

22

CHAPTER 3. ENHANCEMENT OF LUMINESCENCE EFFICIENCY 23

Figure 3.1: As-grown peak PL intensity and emission wavelength of GaIn-NAsSb/GaNAs quantum wells grown under various As-to-Group III flux ratios witha fixed antimony flux [63].

◦C) are required to suppress the phase segregation of InGaNAs alloys with relatively

high concentrations of indium (∼40%) and nitrogen (≥4%) required to achieve 1.55

µm emission [27]. Consequently, near stiochiometric arsenic-to-group III flux ratios

must be used to minimize the introduction of arsenic-related point defects that in-

corporate more readily at these low growth temperatures. Control of the arsenic flux

in this regime is remarkably difficult.

Adding antimony to InGaNAs suppresses phase segregation and therefore higher

substrate temperatures can be used for the growth of GaInNAsSb alloys that emit at

1.55 µm. The arsenic fluxes that can be used at these temperatures are far from the

stiochiometric limit and hence are easier to control. To find the optimal arsenic flux

that yields the highest material quality, an initial study was carried out. The results

of this study were reported in Ref. [63] and are presented here for completeness. In the

initial experiment, a series of GaInNAsSb/GaNAs quantum well samples were grown

at 410 ◦C under various arsenic fluxes. The optical quality of the material was found

to degrade when grown under low arsenic fluxes. Figure 3.1 shows the as-grown peak


0 2 4 6 8 100

0.2

0.4

0.6

0.8

1.0

Arsenic-to-Group III Flux Ratio

Pea

k P

L In

tens

ity (

a. u

.)

410 °C440 °C

Figure 3.2: Peak PL intensity of annealed GaInNAsSb/GaNAs quantum wells grownat 410 ◦C and 440 ◦C under various arsenic-to-group III flux ratios with a fixedarsenic-to-antimony flux ratio. The quantum wells were annealed at 760 ◦C for 60secs.

PL intensity and emission wavelength of the GaInNAsSb/GaNAs quantum wells as

a function of the As-to-Group III flux ratio. For this experiment, the antimony flux

was kept constant and consequently, as the arsenic flux was reduced the As-to-Sb flux

ratio decreased, resulting in higher antimony surface concentrations and increased

antimony incorporation. This was confirmed by secondary ion mass spectroscopy

(SIMS) measurements that show a monotonic increase in the antimony concentration

from 3% at high As/III flux ratios to 12% at low As/III flux ratios [28]. These results

suggest that at low arsenic fluxes, the higher antimony surface concentration reduces

the surface diffusion length below its optimal point—degrading the optical quality

of the material. To evaluate this hypothesis, a set of GaInNAsSb/GaNAs quantum

wells was grown at 410 ◦C under various arsenic fluxes, but using a fixed As-to-Sb

flux ratio. It was found that the PL efficiency of the quantum wells increases with

decreasing group-V fluxes [63].

Motivated by these results, we repeated the experiment using a growth temper-

ature of 440 ◦C, which is the optimal growth temperature for GaInNAsSb/GaNAs


quantum wells emitting at 1.55 µm [64]. For this experiment the antimony flux was

varied along with the arsenic flux to keep a constant and optimal As-to-Sb flux ratio

of 42. In contrast to the initial study, by keeping the As-to-Sb flux ratio constant,

the amount of antimony incorporated in the quantum well remained constant as the

group-V fluxes were varied. This was confirmed by SIMS measurements performed

on selected samples. A significant enhancement in PL efficiency was obtained by

growing at reduced group-V fluxes as shown in Fig. 3.2. The optimal As-to-group

III flux ratio was found to be 4.91, providing a 1.8× improvement over the quantum

well grown at high arsenic fluxes, which was the basis of the lowest threshold lasers.

This enhancement can be attributted to a reduction of the arsenic-related point de-

fects. By using a constant As-to-Sb flux ratio, the antimony surface concentration

and its effects on the surface diffusion length were constant for all samples, and hence

the benefits of using lower arsenic fluxes to minimize point defects were not offset

by the adverse effects of lower surface diffusion lengths. The optical quality of the

GaInNAsSb/GaNAs quantum well degrades if the group-V fluxes are reduced below

the optimal point possibly due to phase segregation and roughening. The sample

grown with an As-to-group III flux ratio of 1.1 was optically inactive and presented

severe structural degradation as is evident in Fig. 3.3 that shows the (004) ω-2θ high-

resolution XRD scans for samples grown under various As-to-group III flux ratios and

substrate temperatures.

In addtion to the quantum well samples grown at 440 ◦C, a series of samples

were grown at 410 ◦C to investigate if there is any additional benefit by growing

at lower temperatures. Lower As fluxes than in the preliminary experiment were

used. At 410 ◦C the optimal group V fluxes were substantially lower than at 440 ◦C

as can be seen in Fig. 3.2. This is because point defects incorporate more readily

at lower growth temperatures and therefore lower arsenic fluxes must be used to

minimize their incorporation. In this case, lower arsenic fluxes were tolerated since

phase segregation is inhibited at lower substrate temperatures. This is supported by

XRD measurements. No discernable structural defects can be seen in the XRD scan

of the quantum well grown at 410 ◦C with an As-to-group III flux ratio of 1.6 and the

1This corresponds to an As/III beam equivalent pressure (BEP) ratio of 9.


-8000 -6000 -4000 -2000 0 2000 4000101

103

105

107

109

ω-2θ (arcsecs)

Inte

nsity

(ar

b. u

nits

)

As/III=1.1 (440°C)As/III=1.6 (410°C)As/III=4.9 (440°C)

Figure 3.3: (004) ω-2θ high-resolution XRD scans of GaInNAsSb/GaNAs quantumwells grown under various arsenic-to-Group III flux ratios and substrate temperatures.

presence of Pendellosung fringes indicate excellent interface quality, in stark contrast

to the quantum well grown at 440 ◦C under similar group-V fluxes (See Fig. 3.3).

The optimal group-V fluxes at 410 ◦C and 440 ◦C yield roughly the same peak PL

intensity as shown in Fig. 3.2. It is preferable to grow at 440 ◦C as the optimal arsenic

flux at 410 ◦C is close to the stiochiometric limit and hence difficult to control.

3.2 Annealing Conditions Optimization

A very high density of defects is introduced during the growth of dilute nitrides. Some

of these defects act as efficient non-radiative recombination centers that degrade the

optical quality of the material. Several types of defects are incorporated, including

gallium interstitials [65], arsenic antisites [66], gallium vacancies [66], and nitrogen

interstitials [67, 68]. The presence of nitrogen on the surface lowers the formation

energy of gallium interstitials [69] causing the introduction of a large concentration

of these defects. Nitrogen split interstitials that incorporate on a single arsenic site

and consits of a nitrogen dimer (N-N)As or a nitrogen-arsenic complex (N-As)Asare also thermodynamically favorable and incorporate readily. Arsenic antisites and


0 720 740 760 780

1520

1540

1560

1580

1600

Ex-situ Annealing Temperature ( °C)

Wav

elen

gth

(nm

)

0 720 740 760 7800

0.2

0.4

0.6

0.8

1.0

Ex-situ Annealing Temperature ( °C)

PL

Inte

nsity

(a.

u.)

a b

Figure 3.4: (a) Peak PL intensity and (b) emission wavelength of GaInNAsSb/GaNAsquantum wells annealed for 60 secs at various temperatures.

gallium vacancies are arsenic-related point defects that incorporate at the low growth

temperatures and high arsenic fluxes required for the growth of dilute nitrides.

Thermal annealing reduces the concentration of defects dramatically improving

the optical quality of dilute nitrides. Arsenic-related point defects and (N-N)As split

interstitials [67, 68] are efficiently removed by annealing whereas gallium interstitials

[65] and (N-N)As split interstitials [67] are more stable and a significant concentration

remains in the material after annealing. Annealing is therefore routinely used to

improve the quality of dilute nitrides.

The standard annealing procedure used in our group was to anneal the quantum

well structures in a rapid thermal annealing furnace for 60 seconds at temperatures

in the 680-820 ◦C range. Figure 3.4 shows the typical annealing behavior for a 1.55

µm GaInNAsSb/GaNAs quantum well. The peak PL intensity increases with anneal-

ing temperature up to an optimal annealing temperature that yields the highest PL

intensity. Above this temperature the PL intensity decreases with increasing anneal-

ing temperature. The degradation of material quality has been ascribed to several

factors including the propagation of arsenic vacancies from the surface [70]. For this

procedure the optimal annealing temperature is between 720 ◦C and 780 ◦C. This


700 720 740 760 780 800 820

0.2

0.4

0.6

0.8

1

Annealing Temperature ( °C), 60 secs.

Pea

k P

L In

tens

ity (

a. u

.)

5 10 15 20 25 30 35Annealing Time (min.), 670 °C

Figure 3.5: Peak PL intensity of GaInNAsSb/GaNAs quantum wells annealed undervarious conditions.

optimal annealing temperature depends on the elemental composition and strain of

the quantum wells, but is independent of the growth temperature [71]. The annealing

behavior of GaInNAsSb quantum wells grown with As2 is indepedent of the arsenic

flux [71], in contrast to InGaNAs quantum wells grown with As4 and annealed in a

Ar/H2 ambient [72]. The blueshift of the emission wavelength upon annealing is due

to a re-arrangement of the nitrogen nearest neighbors, as discussed in section 2.4.

While exploring different annealing conditions, it was found that the optical qual-

ity of GaInNAsSb/GaNAs quantum wells improves substantially by using longer an-

neals at lower temperatures [73]. Figure 3.5 compares the peak PL intensity of two

sets of GaInNAsSb/GaNAs quantum wells: one annealed for 60 seconds at various

temperatures and the other annealed at 670 ◦C for various periods of time. The set

annealed at lower temperatures showed overall higher PL intensities, with the sam-

ple annealed for 25 minutes exhibiting a ∼2× improvement in luminescence efficiency

compared to the sample annealed for 60 seconds at the optimal annealing temperature

of 740 ◦C.

The improvement can be ascribed to the differences in the activation energies for


0 2 4 6 8 10 12

1560

1580

1600

1620

1640

In-situ Annealing Time (hours)

Wav

elen

gth

(nm

)

0 2 4 6 8 10 120

0.2

0.4

0.6

0.8

1.0

In-situ Annealing Time (hours)

Pea

k P

L In

tens

ity (

a. u

.)a b

Figure 3.6: (a) Peak PL intensity and (b) emission wavelength as a function of in-situannealing time for a 1.55 µmGaInNAsSb/GaNAs triple quantum well structure.

the removal of as-grown defects and the formation and propagation of annealing-

induced defects [73]. Since the activation energy is smaller for the former process, at

low temperatures the formation of new defects is delayed more than the removal of

defects.

This annealing procedure is robust as it yields high PL intensities for a wide range

of annealing times. Annealing at lower temperature for longer periods of time also

improves the luminescence efficiency of 1.3 µm GaInNAsSb/GaNAs quantum wells

and lattice matched GaInNAs layers with 1 eV bandgap [73].

The quantum wells experience an in-situ anneal during the growth of the top

Al0.33Ga0.67As cladding layer of edge-emitting lasers or the growth of the top AlAs-

GaAs DBR of VCSELs. The growth of the cladding takes 3 hours while the growth of

the top DBR takes up to 8 hours. To prevent surface roughening, Al0.33Ga0.67As layers

are grown in the 580-620 ◦C range or above 680 ◦C. Growing at high temperatures

ensures minimal incorporation of oxygen in the Al-containing layers. But at this

growth temperature, the quantum well can be overannealed during growth leading to

a degradation of the optical quality of the quantum well material. In fact, lasers grown

at this temperature show no improvement in performance with further ex-situ anneal


0 5 10 15 20 25 30

1500

1520

1540

1560

1580

1600

Ex-situ Annealing Time (min.)

Wav

elen

gth

(nm

)

No In-situ AnnealIn-situ Anneal: 3 hrs

0 5 10 15 20 25 300

0.2

0.4

0.6

0.8

1.0


Pea

k P

L In

tens

ity (

a. u

.)

No In-situ AnnealIn-situ Anneal: 3 hrs

a b

Figure 3.7: (a) Peak PL intensity and (b) emission wavelength of as-grown and in-situannealed (600 ◦C for 3 hours) GaInNAsSb/GaNAs quantum wells as a function ofex-situ annealing time. The ex-situ annealing temperature was 680 ◦C.

[74], indicating possible overannealing. To minimize the risk of overannealing, the

Al0.33Ga0.67As cladding layers were grown at 600 ◦C. But for GaInNAs(Sb) quantum

wells of certain elemental compositions, overannealing can occur even at temperatures

as low as 560 ◦C, leading to high laser thresholds [71].

It is important then to determine if the optical quality of the quantum wells de-

grades during this in-situ annealing. For this purpose, a PL test structure containing

three 7.5 nm GaInNAsSb quantum wells surrounded by 20 nm GaNAs barriers was

grown. The quantum wells were grown at 440 ◦C with an As-to-group III BEP ra-

tio of 9 to obtain the highest luminescence efficiency. After growth, the sample was

cleaved into pieces that were annealed in situ at 600 ◦C for 4 to 13 hours. Arsenic

overpressure was applied during the anneal to prevent arsenic desorption from the

surface. As can be seen in Fig. 3.6, the peak PL intensity of the quantum wells

increases with increasing in-situ annealing time indicating that the active region is

thermally robust, and can withstand in-situ anneals at 600 ◦C for 13 hours or more.

The In content of this quantum well was 40%, but no overannealing was observed for

quantum wells with 39% and 41% In concentrations [75].


1400 1450 1500 1550 1600 16500

0.2

0.4

0.6

0.8

1.0 Reduced As flux 670°C (20 min.)

High As flux 740°C (1 min.)

Wavelength (nm)

PL

Inte

nsity

(a.

u.)

Figure 3.8: PL spectra of GaInNAsSb/GaNAs quantum wells grown and annealedunder various conditions.

To prevent overannealing, the ex-situ annealing conditions must be modified tak-

ing into account the in-situ annealing that occurs due to cladding or mirror layers

grown after the quantum wells. As an example, the optimization of the ex-situ an-

nealing conditions for an edge-emitting laser is presented. For this study, a single

GaInNAsSb/GaNAs quantum well sample was grown. One quarter of the sample

was then annealed in situ at 600 ◦C for 3 hours, emulating the anneal that occurs

during the growth of the top AlGaAs cladding layer. The as-grown and in-situ an-

nealed samples were cleaved into pieces that were annealed at 680 ◦C for several

minutes. Figure 3.7 compares the ex-situ annealing behavior of the as-grown and in-

situ annealed samples. The optimal annealing time for the in-situ annealed sample

is lower than for the as-grown samples. The optimization of the annealing conditions

for VCSELs is presented in Chapter 6.

3.3 Conclusions

The use of reduced group-V fluxes decreases the incorporation of arsenic-related point

defects, enhancing the luminescence efficiency of GaInNAsSb/GaNAs quantum wells.


The optical quality of these quantum wells can be further improved by annealing

them at low temperatures for several minutes to inhibit the formation and propaga-

tion of annealing-induced defects. A 3-4× improvement in luminescence efficiency,

compared to the quantum wells used as the active region for the lowest threshold

GaInNAsSb/GaNAs lasers, is obtained when the optimized group-V fluxes and an-

nealing conditions are combined. This improvement is illustrated in Fig. 3.8 that

compares the PL spectra of a quantum well grown under high group-V fluxes (As/III

flux ratio=10) and annealed for 60 s at 740 ◦C and one grown under reduced group-V

fluxes (As/III flux ratio=5) and annealed at 680 ◦C for 25 min.

The substantial improvement in optical quality is expected to result in lower

laser thresholds. The performance of lasers employing quantum wells with enhanced

luminescence efficiency is presented in the next chapter.

Chapter 4

1550 nm GaInNAsSb Lasers with

GaNAs Barriers

GaInNAsSb/GaNAs quantum wells with enhanced luminescence efficiency were ob-

tained by optimizing the group-V fluxes and annealing conditions. This chapter

presents the development of lasers that use these quantum wells as the gain medium.

The sources of non-radiative recombination and the temperature sensitivity of the

lasers are investigated and strategies to further improve the performance of these

lasers are discussed.

4.1 Laser Structure and Fabrication

Figure 4.1 is a schematic illustration of the structure of the ridge-waveguide lasers

employed in this dissertation. The lasers were grown on n-type GaAs substrates

and followed by a separate confinement heterostructure (SCH), in which the active

region was centered in a waveguide formed by a 460 nm GaAs core region and 1.8 µm

Al0.33Ga0.67As cladding layers. The bottom n-type cladding was doped with silicon

to 7×1017 cm−3 in the half closest to the core region and to 3×1018 cm−3 in the outer

half. The top p-type cladding layer was doped with carbon to 5×1017 cm−3 in the

inner half and to 3×1018 cm−3 in the outer half. The structure was capped by a 50 nm

GaAs layer doped heavily p-type (1×1020 cm−3) to provide low resistance contacts.

33

CHAPTER 4. 1550 nm GaInNAsSb LASERS WITH GaNAs BARRIERS 34

n-GaAs substrate

n-Al0.33

Ga0.67

As

(1.8 µm)

(1.8 µm)

p-Al0.33

Ga0.67

As

Active region

p+-GaAs cap

GaAs core (

Figure 4.1: Schematic of the edge-emitting GaInNAsSb laser structure.

The active region consisted of one or three 7.5 nm GaInNAsSb quantum wells

surrounded by 20 nm strain-compensating GaNAs barriers. The indium, nitrogen and

antimony concentrations in the quantum well were 41%, 2.8%, and 3%, respectively,

and the nitrogen content in the barriers was 3.3%. To obtain the highest luminesence

efficiencies, the quantum wells were grown at 440 ◦C with an As/III flux ratio of 5

and an As/Sb flux ratio of 42. The AlGaAs cladding layers were grown at 600 ◦C.

After growth, the laser structures were ex-situ annealed at 680 ◦C for 10 min-

utes in a rapid thermal annealing furnance using a GaAs proximity cap to minimize

arsenic desorption. These annealing conditions yielded the highest luminesence effi-

ciencies after taking into account the in-situ annealing during the growth of the top

AlGaAs cladding layer (See section 3.2). Laser fabrication consisted of lithography,

Ti/Pt/Au electron-beam evaporation and lift-off to define p-contact stripes of various

widths (5, 10 and 20 µm). The ridge waveguide was then defined using a self-aligned,

chlorine-based electron cyclotron resonance (ECR) plasma etch. Etching the ridge

through the active region reduced slightly the threshold current densities by eliminat-

ing lateral current spreading even at the expense of increased surface recombination.

After etching, the sample was thinned to ∼120 µm and the Au/Ge/Ni/Au n-type


contact was evaporated on the backside. The contacts were annealed at 410 ◦C for 30

seconds. Laser bars of multiple lengths were manually cleaved and soldered epitaxial-

side up on copper carriers using an InSn alloy. These carriers were then mounted on

a temperature controlled copper heatsink using silver paint. Procedures for device

mounting and testing are described in detail in Ref. [76]. No HR or AR coatings were

applied to the laser facets.

4.2 Laser Performance

Figure 4.2 (a) shows the light output versus injected current (L-I) characteristics of

a 20 µm × 2050 µm single GaInNAsSb/GaNAs quantum well laser in continuous

wave (CW) operation at room temperature. The threshold current density was 554

A/cm2, the external differential efficiency was 40%, and the peak output power was

230 mW (from both facets). The lasing wavelength was 1.49 µm as shown in the

optical spectrum in Fig. 4.2 (b). The pulsed (1 µs, 1% duty cycle) L-I characteristics

of the laser are shown in Fig. 4.3. In pulsed operation, the threshold current density

was slightly lower at 540 A/cm2. The maximum output power was 750 mW, limited

by the current source.

The pulsed L-I characteristics of a typical 20 µm × 2450 µm triple quantum well

laser are shown in Fig. 4.4 (a). The threshold current density was 1440 A/cm2, which

corresponds to 480 A/cm2 per well. The external differential efficiency was 25%. The

laser emitted at 1525 nm as shown in the optical spectrum in Fig. 4.4 (b).

Overall, the laser performance was comparable to the performance of the previ-

ous generation of 1.55 µm GaInNAsSb/GaNAs lasers that in CW operation exhibit

threshold current densities as low as 380 A/cm2 for 3 mm-long devices with external

differential efficiencies of 41% and maximum output powers of 250 mW [28]. These

results are suprising since, in general, a significant enhancement in the PL efficiency

of the quantum wells —like the one obtained through the optimization of the group-V

fluxes and annealing conditions— is expected to produce a significant reduction in

laser threshold (see for example Ref. [27]). To understand this result, the sources of

non-radiative recombination in the lasers were investigated as discussed in the next


1470 1480 1490 15000

0.2

0.4

0.6

0.8

1.0

Wavelength (nm)

Pow

er (

a. u

.)

0 200 400 600 800 1000 12000

50

100

150

200

250

Current (mA)

Out

put P

ower

(m

W)

a b

Figure 4.2: Room-temperature, continuous-wave operation of a 20 µm × 2050 µmsingle GaInNAsSb/GaNAs quantum well laser. (a) L-I curve (b) Optical spectrumat thermal rollover.

0 0.5 1.0 1.5 2.0 2.50

200

400

600

800

Current (A)

Out

put P

ower

(m

W)

Figure 4.3: Room-temperature, pulsed (1 µs, 1% duty cycle) operation of a 20 µm ×2050 µm single GaInNAsSb/GaNAs quantum well laser.


1500 1510 1520 1530 15400

0.2

0.4

0.6

0.8

1.0

Wavelength (nm)

Pow

er (

a. u

.)

0 200 400 600 800 1000 12000

10

20

30

40

50

Current (mA)

Out

put P

ower

(m

W)

a b

Figure 4.4: Room-temperature, pulsed (1 µs, 1% duty cycle) operation of a 20 µm× 2450 µm triple GaInNAsSb/GaNAs quantum well laser. (a) L-I curve (b) Opticalspectrum at 1.5× threshold.

section.

4.3 Sources of non-radiative recombination

Several carrier loss mechanisms contribute to the total threshold current density

of a semiconductor laser. These mechanisms include monomolecular, radiative and

Auger recombination. Monomolecular or Schokley-Read-Hall (SRH) recombination

describes the recombination of a single carrier through defects or impurities. In ra-

diative recombination an electron in the conduction band and a hole in the valence

band recombine emitting a photon. Auger is the process whereby an electron-hole

pair recombine across the bandgap, but instead of emitting a photon, tranfers its

energy to a third carrier that is excited higher in energy. This third carrier can be

an electron in the conduction band that is promoted higher in the conduction band

(CHCC process) or a hole in the heavy hole band that is promoted to the split-off

band (CHHS process) [77].

Assuming Boltzmann statistics, the recombination rate of each mechanism is found


to be proportional to the carrier density to the power of the number of carriers involved

in the process. Thus, the current in a semiconductor laser, I, can be expressed as

I = qV (An+Bn2 + Cn3) + Ileak (4.1)

where q is the electronic charge, V is the volume of the active region, n is the carrier

density, and A, B, and C are the monomolecular, radiative, and Auger recombination

coefficients, respectively. Ileak is the leakage current, which is due to carriers that

escape the quantum well and recombine non-radiatively in the barriers or elsewhere

in the laser structure.

All the sources of non-radiative recombination need to be considered when an-

alyzing our 1.55 µm GaInNAsSb/GaNAs lasers. Monomolecular recombination in

the quantum well is possible due to the mid-level traps introduced during growth.

Auger recombination can be significant due to the narrow bandgap (∼0.8 eV) of the

quantum wells. Leakage current is also considerable. As can be seen in the band

diagram for the active region of the lasers plotted in Fig. 4.5, the conduction and

valence band offsets at the GaInNAsSb/GaNAs interface are 104 meV and 109 meV,

respectively. These small band offsets allow carriers to escape thermionically into the

barriers. And since the optical quality of GaNAs alloys with 3% nitrogen is extremely

poor, the carriers recombine non-radiatively in the barriers. This recombination is

exacerbated by the presence of defects, such as gallium interstitials [65] and (As-N)Assplit interstitials [67] that act as efficient electron traps. A significant concentration of

these defects remains in the material after annealing [65,67]. The leakage current into

the GaAs region or the Al0.33Ga0.67As cladding layers is negligible due to the large

conduction and valence band offsets at the GaNAs/GaAs and GaAs/Al0.33Ga0.67As

heterointerfaces.

To determine the relative importance of these non-radiative recombination mech-

anisms, the gain medium of the lasers was characterized by the segmented contact

method [79]. The group of P. Blood at Cardiff University performed the measure-

ments [80]. In this method, the top contact of a broad-area, edge-emitting device is

segmented into multiple electrically-isolated sections of equal length. The amplified


-30 -20 -10 0 10 20 30-0.2

0.0

0.6

0.8

1.0

1.2

∆EC=104 meV

∆Ev=109 meV

GaAs GaAsGaNAs GaNAsGaInNAsSb

E1

E2

HH1HH2HH3

Position (nm)

Ene

rgy

(eV

)

Figure 4.5: Band lineup for the 1.55 µm Ga0.59In0.41N0.028As0.942Sb0.03/GaN0.033As0.967quantum well active region. The diagram was calculated using a conduction bandoffset QC (=∆EC/(∆EC+∆EV )) of 0.85 and 0.8 for the GaInNAsSb/GaAs andGaNAs/GaAs interfaces, respectively [78].

spontaneous emission (ASE) spectrum is then measured when one of the segments is

biased and when two adjacent segments are biased simultaneously. The net modal

gain, internal loss and the spontaneous emission rate spectra are extracted by com-

bining the data from these measurements. From the ratio of gain to spontaneous

emission rate, it is possible to calibrate the spontaneous emission rate in real units.

The radiative current density is obtained by spectrally integrating the calibrated spec-

trum and multiplying by the electronic charge. The non-radiative current density is

obtained simply by substracting the radiative current density from the total injected

current density.

To discern the contribution of barriers and wells to the non-radiative recombina-

tion current, the measurements were performed on both single and triple quantum

well structures. The results are shown in Fig. 4.6 (a) as a function of pumping

level1 [80]. From this data the contribution of wells and barriers to the non-radiative

1Experimentally this is determined as the difference in photon energies at the tranparency pointand at a reference absorption edge. The absorption edge energy was defined as the energy wherethe modal loss was 50 cm−1 for a single quantum well and 150 cm−1 for a triple quantum well.


0 10 20 30 40 50 60 701000

1500

2000

2500

3000

3500

4000

Pumping level (meV)

Non

radi

ativ

e C

urre

nt D

ensi

ty (

A/c

m2 )

Triple QW

Single QW

22 24 26 28 30400

420

440

460

480

500

Pumping level (meV)

Non

radi

ativ

e C

urre

nt D

ensi

ty (

A/c

m2 )

Quantum WellBarrier

a b

Figure 4.6: Segmented contact measurements. (a) Non-radiative recombination cur-rent density as a function of pumping level for single and triple quantum well struc-tures. The open and closed symbols correspond to wells with 3.0% and 3.3% nitrogen,respectively. (b) Contribution of the GaInNAsSb well and GaNAs barrier to the non-radiative recombination current density as a function of pumping level. (Adaptedfrom [80])

current density can be discerned by assuming that the wells and barriers are equally

pumped and that there is no significant leakage into the cladding layers. Three wells

and four barriers contribute to the non-radiative current density of a triple quantum

well structure, JTQW , while one well and two barriers contribute to the non-radiative

current density in a single quantum well structure, JSQW . From these observations the

contribution of well and barrier to the non-radiative current density can be estimated

using the following relations:

Jbarrier =1

2(3JSQW − JTQW ) (4.2)

Jwell = JTQW − 2JSQW . (4.3)

The results are shown in Fig. 4.6 (b) as a function of pumping level. As can be

seen, the contribution of barrier and well to the non-raditive recombination current is


comparable. This means that leakage and subsequent non-radiative recombination in

the barriers is severe and constitutes 66% of the total non-radiative current density

in a single quantum well structure.

The dominant sources of non-radiative recombination in the well and the barriers

were found by matching the experimental results to simulations. The monomolecular

and radiative recombination coefficients (A and B) of the well and barriers used in the

simulations were adjusted to match the measured radiative and total current densi-

ties. The simulations revealed that the main recombination mechanism in both wells

and barriers is monomolecular recombination. As a numerical example, Table 4.1

lists the current densities associated with the various recombination mechanisms at a

pumping level of 25 meV. For a single quantum well structure, the total monomolec-

ular current density is 1220 A/cm2 with the total Auger current density over an

order of magnitude lower at 106 A/cm2. The monomolecular lifetimes in the well and

barrier were found as the inverse of the monomolecular recombination coefficients,

500 and 26 ps, respectively. This large difference reflects the extremely poor optical

quality of the GaNAs (N ∼3%) barriers. The simulations also revealed that Auger

recombination in the barriers is negligible, possibly due to the larger bandgap energy

and the low carrier concentration in the barriers.

These results indicate that the laser thresholds are dominated by leakage into

and monomolecular recombination in the barriers. As most of the non-radiative

recombination occurs in the barriers, a considerable enhancenment of the luminesence

efficiency of the quantum wells produces a marginal reduction of the laser thresholds.

This explains why the performance of the lasers based on the improved quantum

wells is comparable to that of the previous generation of lasers. To further reduce

the thresholds of 1.55 µm GaInNAsSb/GaNAs lasers, it is necessary to improve the

barriers. In addition to affecting the thresholds, leakage into the defect-laden GaNAs

barriers degrades the temperature sensitivity of the lasers as discussed in the next

section.


Recombination mechanism Well (A/cm2) Barrier (A/cm2)

Monomolecular 336 442

Radiative 24.9 0.08

Auger 106 0.02

Table 4.1: Monomolecular, radiative and Auger recombination current densities inthe well and the barriers at a pumping level of 25 meV [80].

4.4 Temperature Sensitivity

The temperature sensitivity of the lasers refers to the variation of the threshold current

density and external differential efficiency as a function of temperature. It was found

empirically that these parameters vary exponentially with temperature as [81]

Jth = J0 exp(T

T0

)(4.4)

ηd = ηd0 exp(− TT1

)(4.5)

where T0 and T1 are the characteristic temperatures of the threshold current density

and external differential efficiency, respectively. These characteristic temperatures

serve as figures of merit for the temperature sensitivity of the lasers. The analysis

of T0 is nontrivial as it depends on the relative importance of each recombination

mechanism along with their temperature dependencies, which are remarkably differ-

ent. The monomolecular, radiative, and Auger recombination coefficients vary with

temperature as follows [82–84]

A ∝ T12

B ∝ T−1

C ∝ exp(−Ea

kT

) (4.6)

where k is the Boltzmann constant and Ea is the activation energy for the Auger

process. This energy is 70 meV for direct processes and falls in the 20-30 meV range

for phonon-assisted processes [83] .


Dominant Mechanism T0 for x ≥ 0 T0 for x = 0

Monomolecular recombination T1.5+x

23T

Radiative recombination T1+2x

T

Auger recombination T3+3x+Ea

kT

≤ 13T

Table 4.2: Characteristic temperature of the threshold current density for ideal(x = 0) and non-ideal (x > 0) quantum well lasers, assuming a dominant recom-bination mechanism. The parameter x accounts for non-idealities that increase thetemperature dependence of the threshold carrier density beyond linearity.

Some insight can be gained by calculating T0 assuming that one of the recom-

bination mechanisms is dominant. For this purpose, it is necessary to consider the

temperature dependence of the threshold carrier density, nth, which in an ideal quan-

tum well laser increases linearly with temperature. Carrier leakage into the barriers

and filling of the higher subbands, however, increase the temperature sensitivity of

nth beyond linearity. To account for these nonidealities nth is expressed as [83]

nth ∝ T 1+x (4.7)

where x is equal to zero for ideal quantum wells.

Relations for T0 when one of the recombination mechanism is dominant are deter-

mined by using the temperature dependencies of the recombination coefficients and

nth in Eq. ( 4.1) and expressing T0 as

1

T0=

d

dTln(Jth). (4.8)

These relations are listed in Table 4.2 for ideal and non-ideal quantum well lasers.

The most temperature sensitive process is Auger recombination leading to T0 values

below 100 K for Auger-dominated lasers. The T0 is higher for lasers where radiative

or monomolecular recombination dominates.

In the case that leakage dominates, the T0 can be estimated from the temperature


0 200 400 600 800 1000 12000

10

20

30

40

Current (mA)

Out

put P

ower

(m

W)

15 °C25 °C35 °C45 °C55 °C65 °C

0 200 400 600 800 1000 12000

10

20

30

40

Current (mA)

Out

put P

ower

(m

W)

15 °C25 °C35 °C45 °C55 °C65 °C

a b

Figure 4.7: Pulsed (1 µs, 1% duty cycle) L-I characteristics as a function of tempera-ture for Ga0.59In0.41N0.028As0.942Sb0.03/GaN0.03As0.97 single (a) and triple (b) quantumwell lasers.

dependence of the thermionic current, which represents a upper bound for the leakage

current. The thermionic current is [85,86]

Je =4πqmikT

2

h3exp

(−Ebi − Efi

kT

)(4.9)

where h is the Planck constant, mi is the effective carrier mass, Ebi is the effective

barrier height and Efi is the carrier quasi-Fermi level. From this relation the T0 can

be calculated as

T0 =T

2 +(Ebi−Efi

kT

) . (4.10)

The pulsed L-I characteristics of the lasers were measured as a function of tem-

perature (See Fig. 4.7) to extract T0, which is equal to the inverse slope of the line

that fits the logarithm of the measured threshold current density versus temperature

data (See Eq.( 4.8)). The measured threshold current density per well as a function of

temperature for single and triple quantum well lasers is shown in Fig. 4.8. The T0 for

single quantum well lasers was 74 K, which means that the threshold current density


280 290 300 310 320 330 3400.2

0.4

0.6

0.8

1.0

1.2

Temperature (K)

Thre

shol

d pe

r w

ell (

kA/c

m2 )

T0=74K

T0=85K

Single Quantum WellTriple Quantum Well

Figure 4.8: Threshold current density per quantum well as a function of temperaturefor single and triple 1.55 µm Ga0.59In0.41N0.028As0.942Sb0.03/GaN0.03As0.97 quantumwell lasers.

roughly doubles when the temperature is increased by 50 ◦C. The T0 of the previous

generation of 1.55 µm single GaInNAsSb/GaNAs quantum well lasers was similar

(71 K) near room temperature, but degraded at elevated temperatures [28]. This

degradation was not observed in the lasers studied here as is evident from Fig. 4.8.

The T0 for triple quantum well lasers was 85 K. The higher T0 can be ascribed to

the fact that in multiple quantum well lasers, some of carriers that escape one well

are captured by another well before recombining.

Since the segmented contact measurements revealed that the main loss mechanism

in the lasers is carrier leakage, the T0 can be estimated using Eq. (4.10). As seen in

Fig. 4.5, the valence band offset at the GaInNAsSb/GaNAs interface is ∼100 meV

and thus the T0 is approximately 16T=50 K. This value is lower than the measured

T0 due to the monomolecular recombination in the quantum wells, which is a less

temperature sensitive process compared to carrier leakage.

The lasers are more temperature sensitive than GaInNAs lasers emitting in the

1.3 µm range that exhibit T0 values above 100 K [87], but less sensitive than 1.55 µm

InP-based GaInAsP lasers, whose T0 is limited by Auger recombination and carrier


leakage to the 40-60 K range [88,89].

Leakage current in the structure can be reduced by improving the barriers. Based

on these results we shifted our attention from the improving the quantum wells to

improving the barriers as discussed in the next section.

4.5 On the improvement of GaNAs barriers

The results of the segmented contact method suggest that to futher reduce the thresh-

old of 1.55 µm GaInNAsSb lasers, it is necessary to reduce the non-radiative recombi-

nation in the GaNAs barriers, or to use an alternative barrier material. This section

describes efforts to improve the GaNAs barriers.

The use of GaNAs barriers was adopted due to several benefits. GaNAs is under

tensile strain on GaAs and hence compensates the high compressive strain (∼2.5%)

of the 1.55 µm GaInNAsSb quantum wells—drastically improving their optical qual-

ity. Additionally, strain compensation enables the growth of multiple quantum wells,

which are required for VCSELs. Furthermore, because GaNAs has a smaller bandgap

compared to GaAs, GaInNAsSb/GaNAs quantum wells are shallower. The smaller

band offsets, although detrimental to the temperature sensitivity of the lasers, result

in lower carrier confinement, allowing longer emission wavelengths with lower indium

and nitrogen concentrations. Figure 4.9 (a) illustrates these advantages by comparing

the PL spectum of annealed GaInNAsSb quantum wells surrounded by GaNAs and

GaAs barriers. The peak PL intensity for the quantum well with GaNAs barriers is

4× higher and its emission wavelength is ∼50 nm longer.

The monomolecular recombination in the barriers can be reduced by decreasing

the barrier thickness. The initial experiment was designed to determine the minimum

GaNAs barrier thickness necessary to mantain adequate optical quality of the GaIn-

NAsSb quantum wells. For this study, a set of 7.5 nm Ga0.60In0.40N0.027As0.938Sb0.035

quantum wells surrounded by GaN0.03As0.97 barriers was grown. With these elemental

compositions, the strain in the well and the barrier is +2.6% and -0.6%, respectively.

The thickness of the barriers was varied from 12.5 nm to 20 nm. After growth, the

samples were annealed for 60 seconds at temperatures in the 700-780 ◦C range. This


1450 1500 1550 1600 1650 17000

0.2

0.4

0.6

0.8

1.0

Wavelength (nm)

PL

Inte

nsity

(a. u

.)

GaNAs Barriers

GaAs Barriers

Figure 4.9: PL spectra of GaInNAsSb quantum wells with GaNAs and GaAs barriers,illustrating the benefits of using GaNAs barriers.

procedure yielded an optimal annealing temperature of 740 ◦C.

Figure 4.10 shows the peak PL intensity of the annealed quantum wells as a

function of barrier thickness. The luminuescence efficiency is similar for quantum

wells with barriers thicker than 15 nm but decreases sharply for quantum wells with

thinner barriers. To understand this behavior, the barrier thickness that results in

zero strain in the structure was calculated. This is simply the thickness of the well

times the strain in the well divided by two times the strain in the barrier. For these

structures, the minimum barrier thickness is 16 nm, sugessting that the strain in the

quantum well must be fully compensated to avoid degradation of the optical quality.

Based on this result, the barrier thickness can be reduced from 20 nm to 16 nm

without degrading the material. Such a small reduction was not expected to produce

a significant improvement in laser performance. This was validated by growing 1.55

µm GaInNAsSb/GaNAs lasers with 16 nm and 20 nm barriers. The threshold current

density for both lasers is comparable as in evident in Fig. 4.11 that shows the L-I

charactersitics of both lasers.

The non-radiative recombination in the barriers can be reduced by improving

the optical quality of GaNAs. This can be accomplished by reducing the nitrogen


12 14 16 18 200

0.2

0.4

0.6

0.8

1.0

GaNAs Barrier Thickness (nm)

Pea

k P

L In

tens

ity (

a. u

.)

Annealing: 740 °C (60 s)

Figure 4.10: Peak PL intensity of 7.5 nm Ga0.60In0.40N0.027As0.938Sb0.035 quantumwells as a function of GaN0.03As0.97 barrier thickness, illustating the importance ofstrain compensation.

concentration, adding dilute amounts of antimony [90] or optimizing the growth con-

ditions [91]. The optical quality of GaNAs is better for alloys with lower nitrogen

concentration, which reduces monomolecular recombination in the barriers. More-

over, the bandgap of GaNAs is wider for lower nitrogen concentrations resulting in

larger band offsets at the GaInNAsSb/GaNAs interface that block carrier leakage

from the quantum well. In practice, this approach is limited by the current configu-

ration of the dilute nitride MBE system used for this work. This system is equipped

with a single gallium source. This forces us to use the same gallium growth rate for

the quantum wells and the barriers, and since the nitrogen concentration is inversely

proportional to the group-III growth rate, the nitrogen content in the barriers is set

by the nitrogen content in the quantum wells. With an additional gallium source,

the nitrogen content in the barrier can be tuned, independently from the nitrogen

content in the wells, to achieve higher quality GaNAs barriers that provide suffficient

strain compensation for the GaInNAsSb quantum wells.

An alternative approach to improve the optical quality of the GaNAs barriers


0 500 1000 1500 2000 25000

20

40

60

80

100

Current Density (A/cm2)

Out

put P

ower

(m

W)

16 nm GaNAs barriers20 µm x 2620 µmJth=1126 A/cm2

20 nm GaNAs barriers20 µm x 2930 µmJth=1085 A/cm2

Figure 4.11: Pulsed (1 µs, 1% duty cycle) L-I characteristics of singleGa0.61In0.39N0.03As0.94Sb0.03/GaN0.033As0.967 quantum well lasers with 16 and 20 nmbarriers.

is to add dilute amounts of antimony to form GaNAsSb alloys with antimony con-

centrations less than 2% [90]. Antimony is supplied by an unvalved cracker in the

MBE system used for this work, which forces us to use the same antimony flux for

the wells and the barriers. The optimal antimony flux for the quantum wells is too

high for the barriers and results in GaNAsSb alloys with antimony concentrations

close to 10% [92], which are nearly lattice-matched to GaAs and hence, provide no

strain compensation for the GaInNAsSb quantum wells. The use of GaNAsSb barriers

with this antimony concentration exarcerbates the leakeage into and monomolecular

recombination in the barriers because, compared to GaNAs, GaNAsSb (Sb∼10%) al-

loys suffer from worse optical qualities and have lower bandgap energies, which result

in smaller band offsets at the well/barrier interface. Morevover, since most of the

bandgap reduction is in the valence band, this can result in type-II band alignment

that provide no hole confinement. The antimony flux can be reduced by repeately

opening and closing the shutter. But attempts to implement this approach were un-

successful [93]. A valved antimony cracker allows rapid modulation of the antimony

flux, enabling the growth of optimal GaInNAsSb quantum wells with high-quality


1400 1450 1500 1550 1600 16500

0.2

0.4

0.6

0.8

1.0

Wavelength (nm)

PL

Inte

nsity

(a.

u.)

Annealing: 740 °C (60 s)

As/Ga = 6.1As/Ga = 12.2

Figure 4.12: PL spectra of annealed 7.5 nm GaInNAsSb quantum wells surroundedby 20 nm GaNAs barriers grown with various As/III flux ratios.

GaNAsSb (Sb≤2%) barriers.

Next, the possibility of improving the quality of GaNAs by modifying the growth

conditions was explored. As the growth temperature is set by the quantum well

growth temperature, the only condition that can be modified is the arsenic flux. To

investigate the effect of the arsenic fluxes, a set of 7.5 nm GaInNAsSb quantum

wells with 20 nm GaNAs barriers was grown. The barriers were grown under various

arsenic fluxes. A slight improvement in luminescence efficiency was observed for

the quantum well with barriers grown at high arsenic fluxes. This is llustrated in

Fig. 4.12 that compares the PL spectrum of quantum wells surrounded by barriers

grown with As/III flux ratios of 6.1 and 12.2. A 10% enhacement in PL intensity

was observed for the sample with barriers grown at an As/III BEP ratio of 12.2.

Dynamical simulations of the (004) ω-2θ XRD scans of the samples revealed slightly

lower nitrogen concentration in the barriers grown under higher arsenic fluxes. Thus,

the improvement in luminescence efficiency can be attributed to the slightly better

optical quality of barriers with reduced nitrogen content. This marginal improvement,

however, is not expected to make a significant impact on laser performance.


4.6 Conclusion

The GaInNAsSb/GaNAs lasers with improved quantum wells showed no improve-

ment over the previous generation of lasers partly because the barriers contribute to

most of the non-radiative recombiantion in the lasers. These barriers are of extremely

poor optical quality and provide insufficent carrier confinement leading to leakage and

subsequent monomolecular recombination. These processes also affect the tempera-

ture sensitivity of the lasers. The laser performance can be improved by reducing the

monomolecular recombination in the barrier or using an alternative barrier material.

The latter strategy is explored in the next chapter.

Chapter 5

1500 nm GaInNAsSb Lasers with

GaAsP Barriers

The previous chapter concluded that GaNAs barriers were limiting the threshold

and temperature sensitivity of the 1.55 µm GaInNAsSb lasers. An alternative barrier

material is therefore required to further improve the performance of these lasers. This

material must be under tensile strain on GaAs to compensate for the high compressive

strain of the GaInNAsSb quantum well as strain compensation is indispensable to

avoid degradation of the optical quality of the quantum wells. In addition, the barriers

need to provide large band offsets with the quantum well to prevent carrier leakage

for improved temperature stability.

GaAsP, InGaP and AlInP alloys meet these requirements as shown in Fig. 5.1.

Among these alloys, GaAsP was selected due to its low growth complexity. GaAsP

has been used succesfully as a barrier material for 1.17 µm InGaAs [94] and 1.3 µm

GaInNAs [95] lasers grown by MOCVD . This chapter presents the growth and char-

acterization of GaInNAsSb quantum wells with GaAsP barriers and the performance

of the lasers that use these quantum wells as the gain medium.

52

CHAPTER 5. 1500 nm GaInNAsSb LASERS WITH GaAsP BARRIERS 53

5.4 5.5 5.6 5.7 5.8 5.9 6.0 6.1

0.5

1.0

1.5

2.0

2.5

Lattice Constant (A)

Ene

rgy

Gap

(eV

)GaAs

AlAsGaP

AlP

GaNAs InAs

InP

GaSb

Figure 5.1: Bandgap energy versus lattice constant for various III-V semiconductors,showing GaAsP, InGaP and AlInP alloys that are under tensile strain on GaAs. Thebandgap energy and lattice constant were calculated using the material parameterscompiled in Ref. [96].

5.1 GaInNAsSb quantum wells with GaAs-GaAsP

barriers

GaAsP layers with ∼18% phosphorous are required to achieve the same degree of

strain as the GaNAs barriers. The incorporation of phosphorous in GaAsP depends

on the substrate temperature and the As-to-P flux ratio [97, 98]. A series of 10

nm GaAsP layers were grown to calibrate the phosphorous flux to achive a desired

phosphorous concentration. One set was grown at 580 ◦C with an As/Ga beam

equivalent pressure (BEP) ratio of 10 and the other set was grown at 440 ◦C with

an As/Ga BEP ratio of 4. The phosphorous concentration in the GaAsP layers was

estimated from the dynamical simulation of the (004) ω-2θ high-resolution XRD scans

of the samples.

Next the GaAsP layers were incorporated into the quantum well active region. As

mentioned in chapter 2 to turn on and off the phosphorous flux, it is necessary to stop

the growth and change the GaP cell temperature. To avoid interrupting the growth


-30 -20 -10 0 10 20 30

-0.2

0.0

0.6

0.8

1.0

1.2

1.4

∆EC=449 meV

∆Ev=178 meV

GaAs GaAsGaAs0.8

P0.2

GaAs0.8

P0.2

GaInNAsSb

E1

E2

E3

HH1HH2HH3

Position (nm)

Ene

rgy

(eV

)

Figure 5.2: Calculated band diagram for a 7.5 nm GaInNAsSb quantum well sur-rounded by GaAs-GaAs0.8P0.2 barriers.

at the quantum well interfaces, a 10 nm GaAs spacer was introduced between the

quantum well and the GaAsP barriers. The calculated band diagram of the structures

is plotted in Fig. 5.2. The conduction and valence band offsets are 462 meV and 172

meV, respectively. These band offsets are higher than the ∼100 meV offsets in the

conduction and valence band provided by the GaNAs barriers.

To investigate the structural and optical properties of this structure, a 7.5 nm

GaInNAsSb quantum well surrounded by 10 nm/20 nm GaAs/GaAsP barriers was

grown. A quantum well with GaNAs barriers was grown in the same run for com-

parison. The GaAsP layers were grown at 580 ◦C with an As/Ga BEP ratio of 10

and an As/P BEP ratio of 1.3. The nominal phosphorous composition in the GaAsP

layers was 20%, which yields approximately the same strain as in the GaNAs barriers.

The structural integrity of the structure was evaluated using x-ray diffraction (XRD).

Figure 5.3 shows the (004) ω-2θ high-resolution XRD scans of the GaInNAsSb quan-

tum wells with GaAs-GaAsP and GaNAs barriers. The structural quality of both

samples is excellent. There are no discernable structural defects and the presence

of Pendellosung fringes along the quantum well features indicates excellent interface

morphology.


-8000 -6000 -4000 -2000 0 2000 4000

102

103

104

105

106

107

ω-2θ (arcsecs)

Inte

nsity

(a.

u)

GaAs-GaAsP BarriersGaNAs Barriers

Figure 5.3: (004) ω-2θ high-resolution XRD scans of GaInNAsSb quantum wells withGaAs-GaAsP and GaNAs barriers.

After growth, the samples were cleaved into pieces that were annealed for 60

seconds at temperatures in the 700-780 ◦C range. The optimal annealing temperature

was 760 ◦C. The PL spectra of the optimally annealed samples is shown in Fig. 5.4.

Compared to the quantum well with GaNAs barriers, the quantum well with GaAs-

GaAsP barriers has 3× lower PL intensity. This is in part due to the lower injection

efficiency of this structure. The emission wavelength of the quantum wells with GaAs-

GaAsP barriers is blueshifted by ∼55 nm. This is expected due to the higher carrier

confinement resulting from the larger band offsets.

5.2 Laser Performance

Ridge-waveguide lasers were grown to evalute the feasibility of using the strain-

compensating GaAsP barriers in the active region of the lasers. The lasers employed

the same structure that was described in Section 4.1. The active region of the lasers

consisted of a 7.5 nm GaInNAsSb single quantum well surrounded by 10 nm/20 nm

GaAs/GaAs0.8P0.2 barriers. Lasers with 20 nm GaN0.03As0.97 barriers were grown in

the same run for comparison.


1400 1450 1500 1550 1600 1650 17000

0.2

0.4

0.6

0.8

1.0

Wavelength (nm)

PL

Inte

nsity

(a.u

.)

GaNAs Barriers

GaAs-GaAsP Barriers

Figure 5.4: PL spectra of GaInNAsSb quantum wells with GaAs-GaAsP and GaNAsbarriers.

After growth, the laser structures were ex-situ annealed at 680 ◦C for 10 minutes.

Figure 5.5 shows the room-temperature, pulsed (1 µs, 1% duty cycle) light output

versus injected current density L-J curves of the lasers. The lasers with GaAs-GaAsP

barriers showed significantly higher threshold current densities than the lasers with

GaNAs barriers, 1.8 kA/cm2 versus 810 A/cm2. The external differential efficiencies

of the lasers with GaNAs and GaAs-GaAsP barriers were 41% and 27%, respectively.

One probable cause for the higher thresholds of the lasers with GaAs-GaAsP bar-

riers is defects introduced during the growth interruptions. In particular, defects can

be introduced at the GaAsP-GaAs interface below the quantum well as the nitrogen

plasma is on during the growth interruption1. Evidence for defects was found in trans-

mission electron micrographs shown in Fig. 5.6. Compared to the top GaAsP-GaAs

interface the bottom interface appears darker. This darker region might be due to a

larger concentration of defects that cause increased scattering of the electron beam.

Figure 5.7 shows secondary ion mass spectroscopy (SIMS) measurements of a

1The nitrogen plasma source was lighted approximately 10 minutes before the growth of thequantum well to allow for plasma stabilization.


0 500 1000 1500 2000 2500 30000

20

40

60

80

100

Current Density (A/cm2)

Out

put P

ower

(m

W)

GaNAs Barriers20 µm x 1860 µmJth=810 A/cm2

GaAs-GaAsP Barriers20 µm x 2410 µmJth=1790 A/cm2

Figure 5.5: Room-temperature, pulsed (1 µs, 1% duty cycle) L-I characteristics ofGaInNAsSb edge-emitting lasers with GaAs-GaAsP and GaNAs barriers.

20 nm

GaAsP

GaAsP

GaAs

GaAs

GaInNAsSb

Figure 5.6: Transmission electron microscopy (TEM) image of a GaInNAsSb quantumwell with GaAs-GaAsP barriers (TEM characterization by Angie C. Lin at StanfordUniversity).


40 60 80 100 1200

1

2

3

4

5

6

Depth (nm)

N, S

b C

once

ntra

tion

(%) In

N

Sb P

5

10

15

20

25

30

In, P

Con

cent

ratio

n (%

)

5

10

15

20

25

30

Figure 5.7: SIMS depth profile for GaInNAsSb quantum wells with GaAs-GaAsPbarriers

GaInNAsSb/GaAs-GaAsP quantum well sample. The measurements showed in-

creased nitrogen concentration at the bottom GaAsP-GaAs interface, possibly due

to nitridation of the interface that occured during the growth interruption. The mea-

surement revealed minimal (below the detection limit) phosphorous incorporation in

the quantum well indicating that the length of the growth interruption is sufficient.

5.3 Temperature Sensitivity

The characteristic temperature of the threshold current density, T0, was extracted

from the dependence of the threshold current density on temperature shown in Fig. 5.8

for both lasers. The T0 for the laser with GaAs-GaAsP barriers was 118 K, signif-

icantly higher than the T0 of 71 K of the laser with GaNAs barriers. However, it

is important to note that this higher T0 might be due to the defects responsible for

the higher thresholds, which provide a monomolecular recombination path that is

temperature insensitive as discussed in the previous chapter.


280 290 300 310 320

1.8

2.0

2.2

2.4

Temperature (K)

Thre

shol

d C

urre

nt D

ensi

ty (

kA/c

m2 )

T0=118K

a b

280 290 300 310 320 330 3400.8

1.0

1.2

1.4

1.6

1.8

Temperature (K)

Thre

shol

d C

urre

nt D

ensi

ty (

kA/c

m2 )

T0=71K

Figure 5.8: Temperature dependence of the threshold current density for GaInNAsSblasers with (a) GaNAs and (b) GaAs-GaAsP barriers.

5.4 Asymmetric GaNAs/GaAs-GaAsP barriers

The previous analysis suggests that most of the defects that cause the higher thresh-

olds are located at the GaAsP/GaAs interface grown before the quantum well. The

incorporation of these defects can be avoided by using a phosphorous valved cracker

that allows rapid modulation of the phosphorous flux, eliminating the need for growth

interruptions.

Since the critical interface appears to be the one below the quantum well, it might

be possible to improve the performance of the lasers by incorporating only the GaAsP

layer that is grown above the quantum well. To investigate this possibility, a set of

GaInNAsSb quantum wells with asymmetric barriers was grown and characterized.

The test structures consisted of a 300 nm GaAs buffer, a 20 nm GaNAs barrier, a

7.5 nm GaInNAsSb quantum well and either a 70 nm GaAs cap layer or a 10 nm

GaAs spacer followed by a 20 nm GaAsP barrier and a 40 nm GaAs cap layer. The

structural quality of both samples is excellent as is evident in the (004) ω-2θ high-

resolution XRD scans shown in Fig. 5.9. After growth, the samples were cleaved

into pieces that were annealed at 680 ◦C for various periods of time. The annealing


-8000 -6000 -4000 -2000 0 2000 4000

102

103

104

105

106

107

ω-2θ (arcsecs)

Inte

nsity

(a.

u.)

GaNAs/GaAsGaNAs/GaAs-GaAsP

Figure 5.9: (004) ω-2θ high-resolution XRD scans of GaInNAsSb quantum wellssurrounded by asymmetric barriers.

behavior is shown in Fig. 5.10. The PL intensity for both samples is comparable,

indicating that the complications with the GaAsP barriers are compensated by the

benefits of strain compensation.

Therefore the use of a GaAsP barrier grown after the quantum well is promising

for the improvement of the lasers. If the laser structure is grown on n-type substrates,

the GaAsP barrier will block the electrons from escaping from the quantum well. On

the other hand, if the structure is grown on p-type substrates, the GaAsP layer will

block the holes from escaping.

5.5 Conclusion

GaAsP barriers were investigated as an alternative barrier material to GaNAs to im-

prove the thresholds and temperature sensitivity of 1.55 µm GaInNAsSb lasers. Lasers

with GaAsP barriers showed improved temperature sensitivity but higher threshold

current densities compared to lasers with GaNAs barriers. The higher thresholds are

probably caused by defects introduced during the growth interruptions required to

change the GaP source temperature. Therefore, lower laser thresholds and improved


0 5 10 15 20 25 300

0.2

0.4

0.6

0.8

1.0

Ex-situ Annealing Time (min)

Pea

k P

L In

tens

ity (

a. u

.)

GaNAs/GaAs BarriersGaNAs/GaAs-GaAsP Barriers

Figure 5.10: Peak PL Intensity as function of ex-situ annealing time for 1.55 µmGaInNAsSb quantum wells surrounded by asymmetric barriers

temperature sensitivity are expected with the use a phosphorous valved cracker that

allows rapid modulation of the phosphorous flux without growth interruptions.

Chapter 6

1550 nm GaInNAsSb VCSELs

The ultimate goal of the development of dilute nitride GaInNAsSb alloys that emit at

1.55 µm is the realization of long-wavelength GaAs-based VCSELs. These VCSELs

are ideal light sources for fiber-to-the-home networks due to their low cost manufac-

turability. The realization of these VCSELs is challenging due to both the difficulties

in the growth of the active region and the complications due to absoprtion that com-

plicate the design of AlAs/GaAs DBRs at 1.55 µm. Despite these challenges, efforts

in our group led to the demonstration of the first electrically-injected GaAs-based

VCSEL emitting beyond 1.31 µm [29]. These VCSELs emitted at 1.46 µm and oper-

ated in pulse mode at -10 ◦C. Similar VCSELs emitting at 1.53 µm that operated in

pulse mode at -40 ◦C were later demonstrated [30].

This chapter presents the design, growth and characterization of the first GaAs-

based 1.55 µm GaInNAsSb VCSELs that operate at room temperature. To improve

the performance of these VCSELs, a new design using two n-doped DBRs and a

tunnel junction is proposed.

6.1 Challenges for 1.55 µm AlAs/GaAs DBRs

While the design of AlAs/GaAs DBRs at 850 and 980 nm is relatively straightfor-

ward there are certain complications that arise for the design of DBRs at longer

wavelengths. These complications are discussed in this section.

62

CHAPTER 6. 1550 nm GaInNAsSb VCSELs 63

0.9 1.0 1.1 1.2 1.3 1.4 1.5 1.6

0.45

0.50

0.55

0.60

Wavelength (µm)

Ref

ract

ive

inde

x co

ntra

st

AlAs-GaAs

0.9 1.0 1.1 1.2 1.3 1.4 1.5 1.60

5

10

15

20

25

30

Wavelength (µm)

Abs

orpt

ion

coef

ficie

nt (

cm-1

)

Babic et al.,JQE 33,1369,1997Henry et al.,JQE 19,947,1983

λ3 relation

a b

Figure 6.1: (a) Wavelength dependence of the absorption of p-type GaAs doped to1×1018 cm−3. The experimental absorption data from literature follow a λ3 relation.(b) Refractive index contrast of the AlAs-GaAs material combination as a functionof wavelength.

The major complication arises from the fact that the optical absorption is sig-

nificantly higher at longer wavelengths. It is therefore more difficult to engineer the

doping profile in the DBR because low doping levels must be used to minimize the

absoprtion loss but this increases the electrical resistance of the DBR and active re-

gion heating. These limitations are particularly severe for p-doped DBRs because

at 1.55 µm, inter-valence-band absorption (IVBA) adds significantly to the total ab-

sorption of p-type materials [99]. As an example, Fig. 6.1 (a) shows the absorption

of p-type GaAs doped to 1018 cm−3 as a function of wavelength. The absorption

coefficient increases from 7 cm−1 at 980 nm to 25 cm−1 at 1550 nm, following a λ3

relation [100]. In addition to higher absorption for p-doped materials, the p-doped

DBRs suffer from higher electrical resistances compared to n-doped DBRs due to low

hole mobilities that lead to low bulk conductivities and large hole effective masses

that limit tunneling and thermionic emission currents across the potential barriers at

the AlAs/GaAs heterointerfaces.

Another difficulty is that the refractive index contrast of the Al(Ga)As-GaAs

material system is reduced at longer wavelengths as shown in Fig. 6.1 (b). With


15 20 25 3099.0

99.2

99.4

99.6

99.8

100

Number of DBR pairs

Ref

lect

ivity

(%

)

980 nm1310 nm1550 nm

15 20 25 3099.0

99.2

99.4

99.6

99.8

100

Number of DBR pairs

Ref

lect

ivity

(%

)

980 nm1310 nm1550 nm

a b

Figure 6.2: Peak reflectivity of (a) undoped and (b) uniformly p-doped (1×1018 cm−3)AlAs/GaAs DBRs as a function of mirror pairs at various wavelengths.

a reduced index constrast more mirror pairs are required to achieve a given peak

reflectivity. In addition, the optical field penetrates deeper into DBRs with lower

refractive index contrast, which increases both absorption and diffraction losses in

the DBRs. An approximate relationship for the effective penetration depth, Lp, is

Lp≈λ

4∆n(6.1)

where λ is the Bragg wavelength and ∆n is the refractive index difference. Clearly,

the smaller the index difference, the longer the penetration depth. Moreover, the

stopband over which the DBR has high reflectivity is reduced with lower index con-

trast.

To illustrate the detrimental effects of higher absorption and lower refractive index

contrast, the peak reflectivity of undoped and uniformly p-doped (1×1018 cm−3)

AlAs/GaAs DBRs was calculated as a function of mirror pairs at 980, 1310 and 1550

nm. Due to the lower index contrast at longer wavelengths, the reflectivities of the

lossless DBRs at 1310 and 1550 nm are lower than at 980 nm for the same number

of mirror pairs as shown in Fig. 6.2 (a). Absorption in the DBRs results in lower

reflectivities for the p-doped DBRs compared to their undoped counterparts. The


high absoprtion at 1550 nm leads to significantly lower reflectivities than at 980 and

1310 nm as is evident in Fig. 6.2 (b). Furthermore, absorption limits the ultimate

(maximum) reflectivity of the DBRs. At 1550 nm, the maximum reflectivity of the

p-doped DBR is 99.65 %. This limited reflectivity precludes the use of p-doped DBR

as the non-emitting mirror, which requires reflectivities over 99.9%.

In addition, the 14λ layers of the DBRs are thicker at longer wavelengths, which

further increases the optical loss and electrical resistance of the DBR stack. For

example, a p-doped AlAs/GaAs DBR with a 99.5% reflectivity at 980 nm requires 16

pairs for a total DBR thickness of 2.5 µm while at 1550 nm it requires 23 pairs for a

total DBR thickness of 6 µm.

The electrical resistance of AlAs/GaAs DBRs designed at 1550 nm is higher than

at 980 nm due to the larger number of pairs required to achieve suitable reflectivities,

the thicker 14λ layers and the limited doping imposed by the higher absorption at long

wavelengths. The higher electrical resistance means that more Joule heating is gen-

erated in the DBR. To further exacerbate the situation, the thicker DBR stacks lead

to higher thermal impedances, which means that the heat generated is dissipated less

efficiently. This self-heating effect limits the output power and maximum operating

temperature and degrades VCSEL performance. This is a critical issue due to the

poor temperature stability of the lasers.

6.2 VCSELs with n- and p-doped DBRs

Despite these challenges GaAs-based 1.53 µm VCSELs operating at room temperature

were demonstated [101]. This section presents the design, growth, fabrication and

characterization of these VCSELs.

6.2.1 Design

A conventional top-emitting structure with p- and n-doped DBRs was employed for

the first demonstration of room temperature operation. For this type of structure,

the bottom n-doped DBR reflectivity was set to 99.95% and the reflectivity of the top


p-doped DBR was designed to achieve a given threshold current density and slope

efficiency. The threshold current density Jth can be calculated using the logarithmic

expression

gth = g0 ln(JthJtr

)(6.2)

where gth is the threshold gain and Jtr and g0 are the transparency current and

material gain of the active region, respectively. The threshold gain is

Γ · gth = αi + αm (6.3)

where Γ is the optical confinement factor, αi is the intrinsic loss and αm is the mirror

loss. The optical confinement factor is defined as the overlap between the active

region and the optical mode. In a VCSEL, it can be expressed as

Γ = ΓtΓl = ΓtΓenhNwdwLeff

(6.4)

where Γt and Γl are the transverse and longitudinal confinement factors, respectively,

Γenh is the gain enhancement factor, Nw is the number of quantum wells, dw is

the thickness of the well, Leff is the effective cavity length, which is equal to the

physical length of the cavity, L, plus the penetration depth in each DBR that can be

approximated by Eq. (6.1). The lateral confinement factor is approximately equal to

1 as the transverse mode profile is defined by the area pumped through the aperture.

The gain enhancement factor can be as high as 2 for thin quantum wells placed at

the maximum of the standing wave. The calculations were performed using a value of

1.8, which is a typical value for VCSELs with multiple quantum well active regions.

The mirror loss is

αm =1

2Leffln(

1

RtRb

). (6.5)

where Rt and Rb are the reflectivies of the top and bottom mirrors, respectively. The

threshold condition can be written as [102]


98.6 98.8 99.0 99.2 99.4 99.6 99.82

3

4

5

6

αi=13.5 cm-1

g0=3180 cm-1

Jtr=1639 A/cm2

Rb,α

=99.95%

24 pairs

Top DBR Reflectivity (%)

Thr

esho

ld C

urre

nt D

ensi

ty (

kA/c

m2 )

Figure 6.3: Calculated VCSEL threshold current density as a function of top DBRreflectivity.

ΓenhNwd

Lgth = αi +

1

2Lln

(1

Rt,αRb,α

)(6.6)

where Rt,α and Rb,α are the reflectivities, including absorption losses, of the top and

bottom DBRs, respectively.

To achieve enough gain, a triple GaInNAsSb/GaNAs quantum well active region

was employed. The transparency current, material gain and internal loss of similar ac-

tive regions were estimated to be 1639 A/cm2, 3200 cm−1 and 13.5 cm−1, respectively,

from cavity length measurements of edge-emitting lasers. Using these parameters, the

threshold current density was calculated as a function of the top mirror reflectivity.

To ensure lasing, the target threshold current density was set to 3 kA/cm2. To obtain

threshold current density below this value, the top p-doped DBR reflectivity must be

higher than 99.5% as can be seen in Fig. 6.3.

The next step in the design is to engineer the grading of the AlAs/GaAs het-

erointerfaces of the DBRs. Abrupt interfaces yield the highest DBR reflectivities but

introduce large discontinuities in the valence and conduction bands due to the differ-

ence in bandgap energies of the high and low index materials. These discontinuities


form potential barriers that hinder carrier transport leading to large series resistances

in the DBRs with a concomitant increase in Joule heating that degrades the VCSEL

performance. This issue is severe for p-doped DBRs because tunneling and thermionic

emission processes are very inefficient due to the large hole effective mass. This prob-

lem can be mitigated by using compositionally graded interfaces that result in lower

potential barriers. In p-doped DBRs, the potential barriers can be completely elim-

inated using a parabolic grading profile and modulation doping [103, 104]. In this

scheme, the doping is modulated across the heterointerface such that the resulting

space charge produces a band bending that compensates the parabolic band bending

created by the compositional grading—resulting in a flat valence band.

These compositional grading profiles are straightforward to implement using metal

organic chemical vapor deposition (MOCVD) as the aluminum concentration can be

easily adjusted by changing the flow rate of the group-III precursors. In contrast,

compositional grading is challenging in MBE because changing the aluminum con-

centration requires a change in the source temperatures. This technique can be used

to implement the grading [105], but it is impractical due to the difficulty of ob-

taining stable and reproducible fluxes after changing the source temperatures during

growth. For this reason, compositional grading is generally implemented by digital

alloying [106]. In this process, a ternary alloy is approximated by a short period

superlattice of thin binary alloys and thus compositional grading is accomplished by

changing the duty cycle of the superlattice. The drawback of this approach is that

it requires a large number of shutter operations, which shortens the lifetime of the

shutter mechanisms.

Alternatively, if multiple group-III sources are available, the grading profile can be

approximated by steps of discrete alloys [107]. This approach was followed using an

MBE system equipped with two gallium and two aluminum sources. For the p-doped

DBR, the AlGaAs/GaAs heterointerfaces were graded in six steps over 30 nm that

approximated the parabolic grading as shown in Fig. 6.4 (a). For the n-type DBR, a

linear grading profile was approximated with 3 steps over 5 nm.

The design of the doping profile is critical as it defines the trade-off between the


0 5 10 15 20 25 300

20

40

60

80

100

Position (nm)

Alu

min

um C

onte

nt (

%)

Step gradingParabolic grading

4 8 12 16 20 24

0

5

10

15

20

25

30

Number of DBR pairs

Ave

rage

Dop

ing

per p

air (

1017

cm

-3)

0 1 2 3 4 5 6Distance (µm)

Fiel

d In

tens

ity (a

. u.)

a b

Figure 6.4: Design of the p-doped Al0.91Ga0.09As/GaAs DBR. (a) Step grading usedto approximate the ideal parabolic grading.(b) Doping profile and electric field as afunction of distance.

optical and electrical properties of the DBR: increasing the doping concentration re-

duces the electrical resistance but increases absoprtion loss and reduces the reflectiv-

ity. Free carrier absorption was reduced by using lower doping concentrations in the

the DBR pairs closer to the cavity where the electric field intensity is higher. For the

p-doped DBR, three average doping levels per pair were used. Figure 6.4 (b) shows

the doping profile of the DBR along with the electric field. The low, medium and

high average doping levels used were 2×1017, 8×1017, and 2×1018 cm−3, respectively.

Similarly, the doping profile of the n-doped DBR was divided in three regions of low,

medium and high doping concentration. To further minimize free carrier absorption,

two doping levels were used in each layer, with the high doping concentration at the

node of the optical field.

With the grading scheme and doping profile defined, the reflectivity of the DBRs

was calculated using the transfer matrix method to determine the number of DBR

pairs necessary to achieve the target reflectivities. The p-doped Al0.91Ga0.09As/GaAs

DBR requires 24 pairs to achieve a peak reflectivity of 99.5%. The n-doped mirror

consisted of a AlAs/GaAs DBR followed by a 4-pair Al0.91Ga0.09As/GaAs DBR. The


total number of pairs necessary to achieve 99.95% was 35.

To confine the current, an oxide aperture defined by the lateral oxidation of AlAs

was designed. The low refractive index oxide aperture confines the optical mode by

index guiding but introduces additional scattering losses. This loss can be substantial,

in particular for small aperture VCSELs, and therefore it must be considered during

the design. An abrupt circular oxide aperture was used for simplicity. For this

type of aperture, the single-pass diffraction loss can be estimated using the fitting

relation [108]

αs = 2.28φ1.190

F 1.3exp {−0.206

φ0F} (6.7)

where F is the Fresnel number and φ0 is the characteristic phase shift of the aperture.

For an aperture of radius a and refractive index n0, the Fresnel number is defined as

F =a2n0

λLc(6.8)

where the equivalent cavity length Lc is equal to the physical length of the cavity

plus the diffraction equivalent distance [109] of each DBR. The characteristic phase

shift of the aperture is

φ0 = (2π

λ)∆ndLc (6.9)

where the distributed index difference ∆nd can be calculated analytically [108] or

estimated using the effective index model [110]

∆nd = nd∆λ

λ(6.10)

where nd is the effective refractive index and ∆λ is the local shift of the resonant

wavelength. The effective refractive index can be approximated by the field-weighted

average refractive index.

The scattering loss depends on the thickness of the aperture and its location in

the DBR stack. Placing the aperture at the node of the electrical field results in

weaker optical confinement but substantially lower diffraction loss [111]. To minimize


2 4 6 8 10

0.05

0.10

0.15

0.20

Aperture Diameter ( µm)

Sca

tterin

g Lo

ss (

%)

30 nm40 nm50 nm

Figure 6.5: Single-pass scattering loss from an abrupt oxide aperture of various thick-nesses placed at the node of the optical field as a function of aperture diameter.

the loss, the aperture was placed at the node of the electric field in the Al0.91Ga0.09As

layer of the top DBR closest to the cavity. The scattering loss for an aperture placed

at this point was calculated as a function of aperture diameter for various aperture

thicknesses. As can be seen in Fig. 6.5, the scattering loss increases with decreasing

aperture diameter. The aperture thickness was selected to be 40 nm. For a 4 µm

aperture device, this aperture adds a 0.2% round-trip loss.

Figure 6.6 depicts schematically the designed VCSEL structure. The triple GaIn-

NAsSb/GaNAs quantum well active region is centered in a 1-λ GaAs cavity. The

bottom n-type mirror consists of a 31-pair AlAs/GaAs DBR followed by a 4-pair

Al0.91Ga0.09As/GaAs DBR. The top p-doped mirror consists of a 24-pair Al0.91Ga0.09As

/GaAs DBR.

6.2.2 MBE growth and fabrication

The epitaxial growth of VCSELs requires precise control of layer thicknesses and

alloy compositions to ensure proper spectral alignment of the cavity resonance and


p-contact (Ti/Pt/Au)

35-pair n-DBR(Al(Ga)As/GaAs)

n-GaAs substraten-contact

(Au/Ge/Ni/Au)

3 GaInNAsSb/GaNAs QWs

24-pair p-DBR(91% AlGaAs/GaAs)

oxide aperture

Figure 6.6: Schematic of the top-emitting 1.53 µm GaInNAsSb VCSEL structure.

the peak of the gain medium. A 1% error in the growth rates leads to a 15 nm gain-

cavity mismatch that results in poor laser performance at best or no lasing at all.

For this reason, the growth rate of each group-III source is calibrated just before the

VCSEL growth. The growth rate of the aluminium (gallium) sources was calibrated

by growing a 1 µm AlAs (GaAs) layer on top of a GaAs (AlAs) layer. The reflectivity

of this structure was measured in situ and matched to simulations to extract the

layer thicknesses. As the fluxes can drift over time, the fluxes were measured and the

source temperature adjusted after each DBR growth.

The VCSEL structures were grown in several steps using two MBE systems that

are connected under UHV. The initial step was the growth of the bottom DBR plus

one quarter of the cavity. The reflectivity of this structure was measured in situ to

determine the DBR peak wavelength. According to this measurement, the VCSEL

structure is redesigned. Next, the remainder of the cavity, including the active region,

is grown in the dilute nitride system. The cavity is intentionally grown 10 nm shorter

than designed to allow for any cavity corrections in case the growth rate was faster

than expected. The appropriate correction is found by measuring the reflectivity of

the partial structure. Finally, the remaining section of the cavity is grown followed

by the top DBR.


31 32 33 34

102

103

104

105

106

107

ω-2θ (Degrees)

Inte

nsity

(a.

u.)

ScanSimulation

Figure 6.7: (004) ω-2θ high-resolution XRD scan of a triple GaInNAsSb/GaNAsactive region.

The oxide aperture was formed using a digitally-alloyed Al0.98Ga0.02As layer that

was integrated into the first Al0.91Ga0.09As layer above the cavity. Digital alloys

are preferable for oxidation layers as their oxidation is more reproducible than that

of uniform Al0.98Ga0.02As layers and they are more structurally stable than AlAs

layers [112].

The active region of the VCSELs consisted of three 7 nm Ga0.59In0.41N0.028As0.946

Sb0.026 quantum wells surrounded by 20 nm strain-compensating GaN0.033As0.967 bar-

riers. These elemental compositions were estimated from the dynamical simulation

of the (004) ω-2θ high-resolution XRD scan of a PL test structure grown during the

VCSEL run (See Figure 6.7). The quantum wells were grown at 440 ◦C with an

As/III flux ratio of 5 (Section 3.1). The AlAs/GaAs DBRs were grown at 590-600◦C.

To optimize the annealing conditions, a quarter of the PL test structure was

subjected to the same in-situ annealing (525 minutes at 590 ◦C) that occurs during

the growth of the top DBR. Another quarter was in-situ annealed at 590 ◦C for 285

minutes. These quarters were then cleaved into pieces that were ex-situ annealed in

a rapid thermal annealing furnace at 670 ◦C for various periods of time. Figure 6.8


0 5 10 15 20 25 300

0.2

0.4

0.6

0.8

1.0


Pea

k P

L In

tens

ity (

a. u

.)

No In-situ AnnealIn-situ Anneal: 285 min.In-situ Anneal: 525 min.

Figure 6.8: Peak PL intensity of as-grown and in-situ annealed GaInNAsSb/GaNAstriple quantum wells as a function of ex-situ annealing time. The in-situ and ex-situannealing temperatures were 590 ◦C and 670 ◦C, respectively.

shows the peak PL intensity of the sample as a function of annealing time. The

samples annealed in-situ showed higher luminescence efficiencies compared to the as-

grown sample, indicating that no overannealing has occured during the growth of the

top DBR. The optimal ex-situ annealing time, however, must be adjusted depending

on the duration of the in-situ annealing. Based on these results and previous annealing

studies, the VCSEL structures were annealed at 670 ◦C for 7 minutes.

The PL spectrum of the sample annealed under optimal conditions is shown in

Figure 6.9 along with the reflectivity spectrum of the VCSEL structure. The quan-

tum well emission wavelength was 1540 nm and the cavity resonance of the VCSEL

structure was at 1530 nm such that the optimal cavity mode-gain peak alignment

occured at 5 ◦C.

VCSEL device fabrication consisted of photolithography, Ti/Pt/Au electron-beam

evaporation and lift-off to form ring p-contacts. Next, mesas with diameters rang-

ing from 18 µm to 36 µm were defined by photolithography and isolated with a

chlorine-based electron cyclotron resonance (ECR) plasma etch. Wet oxidation [113]

at 435 ◦C was then used to form the oxide aperture. After oxidation, the n-contact


1400 1450 1500 1550 1600 1650

0.2

0.4

0.6

0.8

1.0

Wavelength (nm)

Ref

lect

ance

(a.

u.)

1400 1450 1500 1550 1600 1650

PL

Inte

nsity

(a.

u.)

Figure 6.9: Measured reflectivity spectrum of the completed VCSEL structure alongwith the PL spectrum after optimal annealing.

Au/Ge/Ni/Au was evaporated on the backside and annealed at 420 ◦C for 1 min.

6.2.3 Performance

The VCSELs lased at and above room temperature in pulsed mode and near room

temperature in continuous wave (CW) operation [101]. Figure 6.10 (a) shows the

pulsed light output power and voltage versus current (L-I-V ) curves of a device with

a 17 µm diameter current aperture at 25 ◦C. The threshold voltage was 5.6 V and

the threshold current was 18.8 mA, which corresponds to a threshold current density

of 8.3 kA/cm2. The minimum threshold current was observed at 5 ◦C, which is

consistent with the cavity-gain offset. By comparison, 20 µm × 2000 µm as-cleaved

edge-emitting lasers using similar active regions exhibited threshold current densities

in the 1.5-2.0 kA/cm2 range at 25 ◦C. The optical spectrum of the VCSEL measured

at 1.2 times the threshold current is shown in Fig. 6.10 (b), and reveals a lasing

wavelength of 1529.6 nm.

CW operation was observed for small-aperture devices up to 20 ◦C. Figure 6.11

shows the CW L-I curve of a 7 µm diameter current aperture device at different


1526 1528 1530 1532 1534

10-7

10-6

10-5

10-4

Wavelength (nm)

Rel

ativ

e In

tens

ity (

a. u

.)

I=1.2Ith

10 20 30 40 50 60

0.02

0.04

0.06

0.08

0.10

Current (mA)

Out

put P

ower

(m

W)

10 20 30 40 50 60

3

6

9

12

15

Vol

tage

(V

)

a b

Figure 6.10: Pulsed operation of a 17 µm aperture GaInNAsSb VCSEL at 25 ◦C. (a)L-I-V curve. (b) Optical spectrum at 1.2 × threshold.

0 1 2 3 4 50

0.5

1.0

1.5

2.0

Current (mA)

Out

put P

ower

(µ W

)

7 µm apertureCW

-5 °C0 °C5 °C10 °C15 °C20 °C

Figure 6.11: CW L-I curves of a 7 µm aperture GaInNAsSb VCSEL at varioustemperatures.


0 10 20 30 40 50

1

2

3

4

Dissipated power (mW)

Wav

elen

gth

shift

(nm

)

7 µm9 µm11 µm13 µm15 µm17 µm

10 20 30 401528

1529

1530

1531

Temperature ( °C)

Wav

elen

gth

(nm

)

∆λ/∆T=0.098 nm/K

a b

Figure 6.12: Measurements to determine the thermal resistance of the VCSEL struc-ture. (a) Wavelength shift versus dissipated power for various aperture diameters (b)Lasing wavelength as a function of temperature.

temperatures. At 15 ◦C the threshold current was 2.87 mA, which corresponds to a

threshold current density of 7.5 kA/cm2.

The fact that CW operation was not possible at higher temperatures indicates

that there is excessive heating in the device. To determine if heat can be removed

efficiently from the laser, the thermal resistance of the VCSEL structure was mea-

sured. The thermal resistance, Rth, relates the increase in temperature to the increase

in dissipated power

∆T = Rth·∆Pdiss. (6.11)

The thermal resistance can be expressed as

Rth =∆λ/∆Pdiss

∆λ/∆T(6.12)

and thus it can be estimated by measuring the emission wavelength dependence on


dissipated power (∆λ/∆Pdiss) and the emission wavelength dependence on tempera-

ture (∆λ/∆T ). The emission wavelength dependence on dissipated power was mea-

sured for devices of various aperture diameters biased in CW mode below threshold

at a heatsink temperature of 15◦C. Figure 6.12 (a) shows the wavelength shift as a

function of dissipated power. The lasing wavelength dependence on temperature was

measured in pulse mode to avoid Joule heating. As can be seen in Fig. 6.12 (b), the

emission wavelength changes with temperature at an approximate rate of 0.1 nm/K.

The thermal resistance calculated combining these measurements is plotted as a func-

tion of aperture diameter in Fig. 6.13. The thermal resistance of a VCSEL follows

the analytical relation [114]

Rth =1

2ξd(6.13)

that describes the heat flow from a circular heat source of diameter d over an isotropic,

semi-infinite substrate of thermal conductivity ξ. By fitting the experimental data

to this expression the effective thermal conductivity of the structure was estimated

to be 0.36 W/K cm. This value is lower than the average thermal conductivity

of bulk GaAs (0.44 W/K cm) and AlAs (0.91 W/K cm) due to the lower thermal

conductivity of the ternary AlGaAs alloys used in the graded regions of the DBR and

phonon scattering at the GaAs-Al(Ga)As interfaces [115]. The measured thermal

resistance decreases from 1.4 to 0.7 ◦C/mW as the aperture diameter increases from

7 to 17 µm. These relatively low values indicate that heat can be removed efficiently

from the structure. However, the heat generated in the device is high enough to limit

the peak output power and prevent CW operation at higher temperatures. The main

source of heat in the VCSELs is the p-doped DBR, which is very resistive because

the high optical absoprtion of p-doped materials at long wavelengths imposes the use

of low doping concentrations to obtain moderate free-carrier absorption and because

of the non-ideal grading scheme used.

The performance of the VCSELs can be improved by employing a VCSEL struc-

ture without a p-doped DBR, which is the main source of heat. This can be done


6 8 10 12 14 16 18

0.6

0.8

1.0

1.2

1.4

1.6

Aperture diameter ( µm)

Ther

mal

Res

ista

nce

(K/m

W)

Fit to Rth

=(2ξd)-1

Experimental data

Figure 6.13: Thermal resistance of the VCSEL structure as a function of aperturediameter.

by replacing the p-doped DBR with an n-doped DBR, an undoped DBR or a high-

contrast grating (HCG) mirror [116]. Intrinsically, the n-doped DBR is less resistive

than the p-doped DBR because the higher electron mobility and lower electron ef-

fective mass result in lower bulk resistivities and higher tunneling and thermionic

currents at the AlAs/GaAs heterointerfaces. Moreover, higher doping levels can be

used in the n-doped DBR due to the substantially lower free-carrier absorption in

n-doped materials compared with p-doped materials at long wavelengths. The use of

two n-doped DBRs requires a tunnel juntion to provide hole injection into the active

region. This tunnel junction must be optimized to minimize its series resistance.

If an intracavity contacted structure is used, the p-doped DBR can be replaced

with an undoped DBR, which eliminates free-carrier absorption. Since current cannot

be injected through the undoped DBR, hole injection is provided by a p-doped layer

located close to the cavity as shown in Fig. 6.14. This layer must be sufficiently thick

to provide adequate current spreading and efficient hole injection into the active

region. Alternatively, the p-doped DBR can be replaced by a HCG mirror. This

approach has been demonstrated in GaAs-based VCSELs emitting at 850 nm [117,118]

and allows tuning of the emission wavelength [118,119].


undopedDBR

p-type layer

n-DBR

Figure 6.14: Schematic of an intracavity-contacted VCSEL structure, showing thep-doped layer used to provide hole injection into the active region.

In this work, a structure with two n-doped DBRs and a tunnel junction was

selected for ease of fabrication and to avoid the current crowding problems present

in the intracavity contacted structure. The design of the tunnel junction VCSEL is

presented in the next section.

6.3 Tunnel junction VCSELs

As discussed in the previous section, the performance of current 1.55 µm GaInNAsSb

VCSELs can be improved significantly by replacing the highly-resistive p-doped DBR

with a tunnel junction and an n-type DBR. This section presents the development

of GaAs tunnel junctions for this application, as well as the design of a 1.55 µm

GaInNAsSb tunnel junction VCSEL.

6.3.1 GaAs Tunnel Junctions

A tunnel junction is basically a p-n junction in which both the p- and n-regions are

degenerately doped. The tunnel junction is incorporated into the VCSEL structure

such that it is reversed biased when the laser is forward biased. Under these bias


1220 1240 1260 12803.0

3.5

4.0

4.5

5.0

Si Cell Temperature ( °C)

Ele

ctro

n C

once

ntra

tion

(1018

cm

-3)

Figure 6.15: Electron concentration in a 0.5 µm GaAs layer doped with silicon as afunction of the silicon effusion cell temperature.

conditions, electrons in the valence band of the p++ region tunnel into the conduction

band of the n++ region, leaving behind holes that are injected into the active region

where they recombine with electrons injected through the other n-doped DBR.

Highly doped n- and p-type GaAs regions are required to ensure high tunneling

probabilities and consequently low series resistances. P-type carbon doping of GaAs

up to 1020 cm−3 is readily achievable by MBE using a CBr4 source. On the other

hand, for Si-doped GaAs grown under standard conditions, the electron concentration

is limited to ∼5×1018 cm−3. If the silicon concentration is increased above this level,

the additional silicon atoms tend to occupy arsenic sites and associate with silicon

atoms on gallium sites to form electrically neutral Si-Si pairs, thus reducing the ac-

tive electron concentration. This limitation is illustrated in Fig. 6.15 that shows the

electron concentration in a Si-doped GaAs layer as a function of silicon cell temper-

ature. At low growth temperatures, the Si-Si pair formation is suppressed, allowing

higher doping concentrations [120]. This is evident in Fig. 6.16 that shows how the

electron concentration in a Si-doped GaAs layer increases with decreasing substrate

temperature and reaches over 1019 cm−3 for layers grown at 460 ◦C. Higher electron


460 480 500 520 540 560 580

4

6

8

10

12

Growth Temperature ( °C)

Ele

ctro

n C

once

ntra

tion

(1018

cm

-3)

Figure 6.16: Electron concentration in a Si-doped GaAs layer as a function of growthtemperature.

concentrations can be obtained by growing at lower temperatures. However, the low-

est substrate temperature used in this study was 460 ◦C, since there is incomplete

dissociation of the CBr4 molecule at temperatures below 400 ◦C, resulting in reduced

active carbon concentrations and bromine incorporation [121].

To evaluate the effect of the doping concentration of the n++ region on the per-

formance of the GaAs tunnel junctions, a set of test structures was grown and char-

acterized. The tunnel junction test structures were grown on n-type GaAs substrates

and consisted of a 300 nm n-GaAs buffer layer, a 25 nm n++ GaAs layer, a 10 nm

p++ (p=1×1020 cm−3) GaAs layer, a 250 nm p-GaAs layer and a 50 nm p++ GaAs

contact layer. The doping concentration of the n++ layer was varied by growing at

different conditions. In one structure, the n++ layer was grown at 460 ◦C with an

As/III BEP ratio of 10 to obtain a doping concentration of ∼ 1.2×1019 cm−3. As an

example, the simulated band diagram of this tunnel junction test structure is shown

in Fig. 6.17. The n++ layer of the other structure was grown at 580 ◦C with an As/III

BEP ratio of 25 to achieve a doping concentration around 5×1018 cm−3.

Device fabrication consisted of photolithography, Ti/Pt/Au electron-beam evap-

oration and liftoff to form the top contacts. Cylindrical mesas of various diameters


280 300 320 340 360-2.0

-1.0

0.0

1.0

Ec

Ev

p p++ nn++

EF

Depth (nm)

Ene

rgy

(eV

)

Figure 6.17: Energy band diagram for a GaAs tunnel junction test structure with a10 nm p++ region doped to 1020 cm−3 and a 25 n++ region doped to 1.2×1019 cm−3.The n and p regions surrounding the tunnel junction are doped to 1×1018 cm−3.

were then defined by wet etching using the metal contact as the mask. Finally, the

Au/Ge/Ni/Au n-contact was evaporated on the backside and annealed at 410 ◦C for

30 seconds.

The forward bias current density-voltage (J-V ) characterisitics of the tunnel junc-

tion with the n++ layer grown at 460 ◦C are shown in Fig. 6.18 (a). The tunnel

junctions exhibit clear tunneling behavior with a peak current density of 245 A/cm2

and a peak voltage of 200 mV. The tunnel junctions with the n++ layer grown at

580 ◦C have peak current densities of 24 A/cm2 as shown in Fig. 6.19 (a). In the

VCSEL structure, the tunnel junction is reverse biased and the voltage drop at the

operating current is an important figure of merit. Fig. 6.18 (b) shows the reverse-

biased IV characteristics of the tunnel junctions with the n++ layer grown at 460 ◦C.

At 5 kA/cm2 the voltage drop is 0.76 V for the as-grown tunnel junction. This value

increases to 1.10 V after annealing the tunnel junction at 760 ◦ for 60 seconds. For

the tunnel junctions with the n++ region grown at 580 ◦C, the voltage drop is 1.33 V

and increases to 1.6 V after annealing as shown in Fig. 6.19 (b).

The voltage drop in the n-doped DBR is estimated to be 0.5 V making the total


0 1 2 3 4 50

0.2

0.4

0.6

0.8

1.0

1.2

Vol

tage

(V

)

Current Density (kA/cm2)

As-grownAnnealed

0 0.2 0.4 0.6 0.80

50

100

150

200

250

Voltage (V)

Cur

rent

Den

sity

(A

/cm

2 )a b

Figure 6.18: Measured (a) forward-biased and (b) reversed-biased current-voltagecharacteristics of a GaAs tunnel junction with ∼1.2×1019 cm−3 active donors in then++ layer.

voltage drop in the top DBR 1.6 V. This is substantially lower than the voltage drop

in the p-doped DBR, which is roughly 4 V. The lower resistance in the top DBR

reduces the Joule heating and allows CW operation at higher temperatures.

It is worth noting that even better performance can be achieved with the use of

other GaAs-compatible tunnel junction material such as dilute nitride p+-InGaAs/n+-

GaInNAs [122] and type-II p+-GaAsSb/n+-InGaAs [123].

6.3.2 Design

The tunnel junction/double n-doped DBR VCSEL design methodology followed was

similar to that used for the design of the VCSELs with p- and n-doped DBRs. An

additional step in the design was the placement of the tunnel junction in the top

DBR. The tunnel junction must be located above the oxide aperture and, to min-

imize absorption, it must be placed at the node of the electric field. This can be

accomplished by introducing a 34λ GaAs spacer after the Al0.91Ga0.09As layer nearest

to the optical cavity and placing the tunnel junction at the node of the electric field

in the spacer as depicted in Fig. 6.20.


0 1 2 3 4 50

0.5

1.0

1.5

2.0

Vol

tage

(V

)

Current Density (kA/cm2)

As-grownAnnealed

0 0.2 0.4 0.6 0.80

5

10

15

20

25

Voltage (V)

Cur

rent

Den

sity

(A/c

m2 )

a b

Figure 6.19: Measured (a) forward-biased and (b) reversed-biased current-voltagecharacteristics of a GaAs tunnel junction with ∼4.5×1018 cm−3 active donors in then++ layer.

4.5 5.0 5.5 6.0 6.5 7.0 7.52.8

3.0

3.2

3.4

Position (µm)

Inde

x of

Ref

ract

ion

(3/4) λ GaAs Spacer Quantum Wells

Tunnel JunctionOxide Aperture

Figure 6.20: Refractive index profile and electric field intensity in the designed tunnel-juction VCSEL structure, showing the location of the tunnel junction at the node ofthe electrical field in a 3

4λ GaAs spacer in the top n-doped DBR.


n-contact (Au/Ge/Ni/Au)

35-pair n-DBR(Al(Ga)As/GaAs)

n-GaAs substraten-contact

(Au/Ge/Ni/Au)

3 GaInNAsSb/GaNAs QWs

22-pair n-DBR(91% AlGaAs/GaAs)

oxide aperture

GaAs tunnel junction

Figure 6.21: Schematic of the designed 1.55 µm GaInNAsSb tunnel-junction VCSELstructure.

To achieve a threshold current density of ∼3 kA/cm2, the reflectivities of the bot-

tom and top DBRs were set to 99.95% and 99.5%, respectively. The Al(Ga)As/GaAs

heterointerfaces of the n-doped DBRs were graded in 3 steps over 4 nm approximating

a linear grading between GaAs and Al0.36Ga0.64As. To reduced free-carrier absorp-

tion, the three average doping levels across the DBR were used. The low, medium

and high average doping concentrations per DBR pair were 4×1017, 7.5×1017, and

2.5×1018 cm−3, respectively.

The designed tunnel-junction VCSEL structure is shown schematically in Fig. 6.21.

The bottom mirror consisted of a 31-pair AlAs/GaAs DBR followed by a 4-pair

Al0.91Ga0.09As/GaAs DBR for a total of 35 pairs. The top mirror is composed of a

22-pair Al0.91Ga0.09As/GaAs n-doped DBR. In contrast, 24 pairs are necessary for a

p-doped DBR to achieve 99.5% reflectivity. The lower number of pairs is mainly due

to the fact that n-doped materials have lower free-carrier absoprtion than p-doped

materials at 1.55 µm. The GaInNAsSb/GaNAs triple quantum well active region is

centered in a 1-λ GaAs cavity.

The VCSEL structures have been grown and the fabrication is underway. Fig-

ure 6.22 shows the in-situ reflectivity of the structure with a cavity resonant wave-

length of 1530 nm.


1400 1500 1600 17000

0.2

0.4

0.6

0.8

1

Wavelength (nm)

Ref

lect

ivity

(a.

u)

Figure 6.22: Measured reflectivity spectrum of a GaInNAsSb tunnel junction VCSEL.

Chapter 7

Conclusions

This thesis presented important steps towards the realization of low-cost GaAs-based

1.55 µm GaInNAsSb lasers for fiber-to-the-home networks.

The optical quality of 1.55 µm GaInNAsSb quantum wells surrounded by strain-

compensating GaNAs was enhanced substantially by optimizing the growth and an-

nealing conditions. The use of reduced group-V fluxes decreased the incorporation of

arsenic-related point defects and led to a 2× improvement in luminescence efficiency.

An additional 2× improvement was obtained using relatively low temperature, long

anneals that inhibit the formation and propagation of annealing-induced defects.

Despite this enhancement, edge-emitting lasers based on the improved quantum

wells showed no reduction in laser thresholds. By using the segmented contact

method, the sources of non-radiative recombination in the lasers were identified.

These measurements revealed that the dominant carrier loss mechanism is leakage

into, and subsequent monomolecular recombination in, the defect-laden GaNAs bar-

riers. This mechanism is also responsible for the poor temperature sensitivity of the

lasers. Approaches to improve the quality of the GaNAs barriers—such as reducing

the nitrogen concentration and adding dilute amounts of antimony—were discussed

but were not implemented due to limitations of the MBE systems used for this work.

GaAsP alloys were explored as an alternative barrier material to reduce the laser

thresholds and improve the temperature sensitivity of the lasers. GaInNAsSb lasers

with GaAsP barriers showed improved temperature sensitivity but higher threshold

88

CHAPTER 7. CONCLUSIONS 89

current densities. The higher thresholds are likely caused by defects introduced during

the growth interruptions required to change the temperature of the GaP source used

to supply phosphorous. The use of a phosphorous valved cracker eliminates the need

for growth interruptions and is expected to enable the realization of low threshold

lasers with improved temperature stability.

This thesis also presented the realization of the first electrical-injected GaAs-based

VCSELs emitting in the 1.55 µm range that operate at and above room temperature

in pulsed mode and near room temperature in continuous-wave mode. The main

factor limiting VCSEL performance and preventing continuous-wave operation at

higher temperatures is Joule heating generated in the highly resistive p-doped DBR.

The problem can be alleviated by replacing the p-doped DBR with another n-doped

DBR, which has lower electrical resistance and lower absoprtion loss. In this VCSEL

structure, a tunnel junction is employed to provide hole injection into the active

region. GaAs tunnel junctions were developed for this application and a tunnel-

junction 1.55 µm GaInNAsSb VCSEL was desgined.

These tunnel-junction VCSELs are expected to lase in continuous-wave mode at

room temperature. The ultimate goal of realizing low-cost, 1.55 µm GaAs-based

VCSEL for fiber-to-the-home networks might be within reach.

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